Structure and Dynamics of Biobased Polyester Nanocomposites

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Structure and Dynamics of Biobased Polyester Nanocomposites Krystalenia Androulaki, Kiriaki Chrissopoulou, Daniele Prevosto, Massimiliano Labardi, and Spiros H Anastasiadis Biomacromolecules, Just Accepted Manuscript • DOI: 10.1021/acs.biomac.8b01231 • Publication Date (Web): 28 Nov 2018 Downloaded from http://pubs.acs.org on November 30, 2018

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Biomacromolecules

Structure and Dynamics of Biobased Polyester Nanocomposites Krystalenia Androulaki,1,2 Kiriaki Chrissopoulou,1,* Daniele Prevosto,3 Massimiliano Labardi3 and Spiros H. Anastasiadis1,2, 1

Institute of Electronic Structure and Laser, Foundation for Research and Technology Hellas, P. O. Box 1527, 711 10 Heraklion Crete, Greece

2

Department of Chemistry, University of Crete, P. O. Box 2208, 710 03 Heraklion Crete, Greece 3

CNR-IPCF, Department of Physics, University of Pisa, Pisa, Italy

Keywords: Bio-based polyesters, Layered silicates, Intercalation, Dynamics, Dielectric Relaxation Spectroscopy



Corresponding Authors. E-mails: SHA : [email protected], KC : [email protected]

ACS Paragon Plus Environment

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Abstract

The structure and the dynamics of two bio-based polyester polyols are investigated in the bulk and close to surfaces in polymer / layered silicate nanocomposites. The morphology of the neat polymers as well as the structure of the nanohybrids are investigated with X-ray diffraction and their thermal properties are studied by differential scanning calorimetry. One of the investigated polyesters is amorphous whereas the second one is a semi-crystalline polymer with an intriguing thermal behavior. Hybrids have been synthesized over a broad range of compositions and intercalated structures are always obtained. The thermal transitions in the nanocomposites are observed only when the polymers are in excess outside the completely filled galleries. The glass transition, whenever it can be resolved, appears insensitive to the presence of the inorganic material whereas the way the crystallization takes place depends on the composition of the nanohybrid. Dielectric relaxation spectroscopy was utilized to study the polymer dynamics. It revealed multiple relaxation processes for the neat polymers both below and above their glass transition temperatures whereas, in the nanocomposites, similarities and differences are observed depending on the specific mode of dynamic process.


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I.

Introduction

Biobased and biodegradable polymers have received much attention over the last three decades and a number of industries tend to substitute petroleum-based with bio-based intermediate materials in order to reduce the environmental footprint of their products.1,2,3,4,5 Among such materials, bio-based aliphatic polyesters have attracted considerable attention due to the combination of biodegradability, biocompatibility and their physicochemical properties, which, in many cases, are comparable to those of more common and widely utilized polymers.6,7,8,9,10,11 Their significant biodegradability is crucial for the reduction of the so-called “white pollution” to be enforced during the following decades whereas their biocompatibility renders them suitable for biomedical applications as well.12,13 In these materials, the correlation between the chemical composition and the final properties is very complex; for example, the number of methylene groups in their repeating unit and the chain length may affect the thermal properties, whereas they may induce important differences in the crystalline structure and morphology.11,14 One of the most important biodegradable polyesters is poly (ethylene succinate), PESu, which is commercially available as well. It is a semicrystalline aliphatic polyester that is easily processable from the melt, is resistant to chemicals and exhibits mechanical properties similar to widely used conventional polymers like polyethylene.15 The main property of interest for biodegradable polymers is their rate of biodegradation, which can be influenced by many factors like macromolecular characteristics (chemical structure, stereoregularity, molecular weight and molecular weight distribution), surface properties (surface area, hydrophilicity / hydrophobicity) as well as thermal and mechanical properties (glass transition temperature, melting temperature, modulus of elasticity). Moreover, the crystalline structure, the total crystallinity as well as the sizes of the spherulites and the lamellae significantly



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affect not only the material properties but the mechanism and the rate of biodegradation as well.16 Crystallinity is generally accepted to hinder biodegradability, which starts from the amorphous regions with much higher erosion rates than those of the crystalline regions. Subsequently, during biodegradation, holes are formed in the center of the spherulites with their size increasing with time; this indicates that the spherulite centers are composed of less ordered lamellae that can degrade more easily.17 One way to influence the surface properties as well as the crystallinity and crystalline properties of a polymeric material is by blending it with another polymer18,19,20,21 or by the incorporation of nanometric additives like hydrophilic or organophilic layered silicates,22 silica nanoparticles,23 graphene derivatives24 that will be dispersed in the polymer matrix forming a nanocomposite. Polymer nanocomposites, comprised of a polymer matrix and inorganic or carbon additives (e.g., nanoparticles, clays, nanotubes, graphene, etc.) as the nanofiller, possess improved and often innovative physicochemical properties compared to conventionally filled systems.25,26,27,28,29,30,31 These superior properties are attributed to the nanometric filler size that leads to a dramatic increase in the interfacial area, thus creating a large volume fraction of “interfacial” polymer chains with properties different from those in the bulk even at small loadings.32,33 Polymer / layered silicate nanocomposites have been widely investigated34,35,36 due to the improvement of mechanical, barrier, flammability, optoelectronic and other properties important for various applications such as membranes for fuel cells or separation devices, food packaging, photovoltaic devices, chemical or biochemical sensors, etc. In such nanocomposites, three types of structures have been observed: the phase separated, the intercalated and the exfoliated one; the final structure is controlled by the interactions between the polymer and the inorganic material.37,38,39 In the intercalated systems, a 0.8-2.5 nm polymer film resides within the galleries formed by the adjacent


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parallel inorganic layers forming a well-ordered multilayer with a repeat distance of a few nanometers. These hybrids are scientifically interesting for the extra reason that it allows the investigation of the structure, conformation and dynamics of polymers under confinement utilizing, however, conventional analytical techniques and macroscopic samples. Moreover, the way polymers crystallize under confinement and/or close to surfaces and how one can alter their conformations or control their crystallization is a fundamental problem and, at the same time, it is of significant importance for technological applications.40,41,42,43,44 The layered silicate nanoparticles are usually hydrophilic and, thus, hydrophilic polymers interact favorably with the silicate surfaces leading mostly to intercalated nanocomposites.42,43 On the other hand, hydrophobic polymers can lead to intercalated45,46 or exfoliated47,48 structures when mixed with organophilized clays, i.e., with materials where a cation exchange reaction is utilized to replace the hydrated cations within the galleries by proper cationic surfactants (e.g., alkylammonium). For polyolefins, which are even less polar polymers, appropriate compatibilizers are required in order to achieve exfoliation by controlling the interactions.37,38,39 Nanohybrid materials composed of biodegradable polymers and different additives (generally in small concentrations) have attracted the scientific interest over the recent years, as well.22,23,24,49,50,51,52,53,54,55,56 Nevertheless, despite the fact that there have been reports on aliphatic polyester nanocomposites with improved properties for biomedical applications in tissue engineering57 as well as nanocomposites with significantly improved mechanical and thermal properties,58,59 research work is still necessary to understand the structure-properties relationships in these systems. An area that has attracted less attention from the scientific community regarding biobased polymers and their nanocomposites is the investigation of their dynamics although understanding of their relaxation processes is crucial in order to derive the optimum processing conditions and,



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thus, expand the application range of such materials.60,61,62,63,64 In general, the investigation of the dynamics of polymers involves the study of bond vibrations, rotational motion of the different side groups, the -process or segmental relaxation and the motion of the whole chain; it is, thus, clear that an extended temporal range of more than ten orders of magnitude, from pico-seconds (ps) to seconds (s), should be covered. The relaxation processes that are observed below the polymer glass transition temperature like the local rotation and reorientation of side groups, usually exhibit an Arrhenius temperature dependence,  = 0 exp[E/RT]. The segmental dynamics or alpha (-) relaxation process, which is due to the cooperative motion of chain segments, is probed above the polymer glass transition temperature, Tg. Its relaxation times do not follow the Arrhenius equation as the temperature approaches the Tg and the behavior is analysed with the Vogel-FulcherTammann equation,   exp  B / T  T0   . The investigation of polymer dynamics under confinement, close to interfaces or in thin films has attracted the interest of the scientific community since significant differences emerge when the molecules are confined over distances comparable to their sizes.65 The investigation of polymer dynamics in intercalated polymer / layered silicate nanocomposites provides the opportunity to study the influence of both confinement and of the different polymer / surface interactions utilizing macroscopic techniques. A number of investigations showed66,67,68,69,70 that the segmental relaxation process of severely confined polymers is much faster than the respective one of the pure polymer, it is observed even at temperatures below the bulk polymer Tg and follows an almost Arrhenius temperature dependence. Nevertheless, in cases where there are more attractive polymer-surface interactions, the probed segmental relaxation shows much slower dynamics than that of the unconfined polymer.71,72,73,74 Moreover, local sub-Tg motions show similar relaxation times67 between the


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confined polymer and the bulk except from the cases that restrictions in motion are imposed by, e.g., the formation of hydrogen bonds, where the local processes under confinement are found faster and with lower activation energies than the respective ones of the neat polymers.75 In the present work, the structure, thermal properties and dynamics of two biobased polyester polyols, one amorphous and one semi-crystalline, are investigated both in the bulk and, for the first time, under confinement within the galleries of sodium montmorillonite, Na+-MMT, in intercalated nanocomposites, which were synthesized over a large range of compositions utilizing solution mixing in water. X-ray diffraction was utilized to study the structure of the polymers and of the nanocomposites. Intercalated nanohybrids having mono- and bi-layers of polymer within the galleries of the inorganic material are obtained in all cases. Investigation of the thermal properties by differential scanning calorimetry reveals that the glass transition, whenever it can be resolved, seems insensitive to the presence of the inorganic material, whereas the way the crystallization proceeds depends on the composition of the nanohybrid. Dielectric relaxation spectroscopy was utilized to investigate polymer dynamics in the bulk and when almost all chains are severely confined within the galleries of the inorganic material, i.e., for a nanocomposite with ~30wt% polymer; the segmental and local sub-Tg relaxation processes are studied for temperatures both above and below the bulk polymer glass transition temperature. The -relaxation process reflects the difference in the glass transition temperature and becomes slower for both polymers in the nanocomposites retaining the VFT temperature dependence. The local dynamics of the processes probed below Tg shows an interesting behavior displaying both similarities and differences with the respective one of the neat polymers. The investigation of the structure and dynamics of bio-based polymers is very important in order to understand their physicochemical properties as well as their biodegradability and



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biocompatibility behavior. Reduced crystallinity, for example, is critical when biodegradation is concerned since crystallinity hinders degradation that begins at the amorphous regions. At the same time, knowledge on the dynamics of these polymer over multiple length and time scales is critical, as well, in order to optimize processing conditions.

Introduction of inorganic

nanoadditives in bio-based polymers is frequently utilized to improve their mechanical properties. Thus, studying how polymer structure and dynamics is affected by the presence of inorganic surfaces and/or confinement in polymer nanocomposites is very significant as well.

II.

Experimental Section

Materials Two different commercial bio-based polyester polyols (Myriant) made from bio-succinic acid are utilized in this work. Poly (ethylene succinate), EG-110, and poly(diethylene succinate), DG110, had been synthesized by the polycondensation reaction of succinic acid with ethylene glycol and diethylene glycol, respectively. Succinic acid that is abundant in nature in plant and animal tissues has the same chemical properties as adipic acid, which is regularly utilized for the synthesis of polyester polyols and, thus, it can generally be introduced without changes to the manufacturing process to increase to a large extend the renewable content.76 The chemical formulae of the two polymers are shown in Scheme 1. Both polymers contain hydroxyl end-groups whereas the repeat unit of DG-110 contains an additional ether linkage. EG-110 is a crystalline polymer whereas DG-110 is an amorphous one, as will be discussed later on. The molecular characteristics of the polymers are given in Table 1 according to the material data sheets as provided by the supplier. Moreover, 1H-NMR


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measurements of the two oligomers have been performed (not shown), which, by proper integration of the NMR peaks, reveal that EG-110 is a polyester composed of six (6) monomers and two end groups, whereas DG-110 comprises of 4-5 monomers and two end groups, one of which is diethylene glycol.

(a)

(b)

Scheme 1. Chemical formulae of (a) poly(ethylene succinate), EG-110, and (b) poly(diethylene succinate), DG-110. Table 1. Characteristics# of bio-based polyester polyols

Polymer

Commercial name

OHnumber*

Acid value⁑

Functionality

Bio-based carbon content

Molecular

(%)

(g/mol)

weight ⁂

Poly(ethylene succinate)

EG-110

107

0.52

2

66

1043

Poly(diethylene succinate)

DG-110

113

0.8

2

50

986

# Table contents according to the material data sheets as provided by the supplier * The hydroxyl value is a measure of the content of free hydroxyl groups in a chemical substance. It is defined as the number of milligrams of potassium hydroxide (KOH) required to neutralize the acetic acid taken up on acetylation of one gram of the substance that contains free hydroxyl groups. ⁑ Acid value is defined as the mass of potassium hydroxide (KOH) in milligrams that is required to neutralize one gram of the chemical substance. ⁂ Mw calculated by Mw = [54.1  1,000 / (OH number + acid number)]  functionality, according to the data provided by the supplier



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The multilayer sodium montmorillonite, Na+-MMT, Cloisite Na+, supplied by Southern Clay, was used in this investigation as the inorganic material for the synthesis of the nanocomposites. Sodium montmorillonite is a layered silicate clay, which is hydrophilic due to the hydrated sodium cations that exist within its galleries. Due to this hydrophilicity, Na+-MMT can be mixed with hydrophilic polymers such as the polyester polyols utilized in this work without the need of any interfacial modification. The period of the multilayer structure is ~1 nm (when completely dry) and exhibits a cation exchange capacity, CEC, of 92.6 mmol/100g. Nanocomposites were prepared via a solution-intercalation method using deionized water as the solvent; the polymers were mixed with Na+-MMT in water to prepare hybrids over a broad range of compositions. The required amount of polyesters was first dissolved in water at room temperature, and, then, the respective amount of clay was added under continuous stirring in order to ensure that the individual clay layers were well dispersed. Following mixing, the solvent was evaporated initially in open air until a concentrated slurry was obtained, which was, then, transferred to a vacuum oven to completely dry. Finally, all samples were thermally annealed at 130C overnight under vacuum to erase any metastable structures formed during solvent evaporation and to achieve equilibrium. Thermogravimetric Analysis (TGA) was utilized to verify the composition of the synthesized hybrids.

Experimental Techniques X-ray Diffraction (XRD). The structural characterization of the two polymers and the respective nanocomposites was performed by X-ray diffraction utilizing a RINT-2000 Rigaku Diffractometer. The X-rays are produced by a 12 kW rotating anode generator with a Cu anode and the Cu Kα radiation with wavelength  = CuKα = 1.54 Å is used. Measurements were


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performed for diffraction angles, 2, from 1.5° to 30° with step of 0.02°. The X-ray diffractograms of the layered silicates show the characteristic (00l) peaks due to their periodic structure. These diffraction peaks are related to the spacing of the layers according to Bragg’s law, nλ = 2d00l sin, where  is the wavelength of the radiation, d00l the interlayer distance, n the order of diffraction, and 2 the diffraction angle. Polymer intercalation leads to a shift of the main peak of the pure inorganic clay to lower angles whereas no peaks are observed in the case of exfoliated structures. Differential Scanning Calorimetry (DSC). The thermal properties of the two polymers and the respective nanocomposites were measured with a PL-DSC (Polymer Laboratories) Differential Scanning Calorimeter, DSC. The temperature range from −100 to 150 °C was covered with a heating / cooling rate of 10 °C/min. In all cases, the transition temperatures, either the glass transition temperature, Tg, for DG-110 or the melting temperature, Tm, for EG-110, were obtained from the second heating cycle in order to eliminate any thermal history and to remove any remaining humidity. The measurements were performed under constant nitrogen gas flow to prevent degradation of the samples and controlled cooling was achieved using liquid nitrogen. Dielectric Relaxation Spectroscopy (DRS). Dielectric spectroscopy measures the dielectric impedance of a medium as a function of frequency. It is based on the interaction of an external field with the molecular electric dipole moment and the determined quantity is the complex dielectric permittivity: *() = '() – i "()

(1)

where ' represents the real part and " the imaginary part or the energy loss part,  = 2π and  the measured frequency. The measured * is given by the one-sided Fourier transform of the time derivative of the dipole-dipole correlation function C(t), which probes local motions unless there



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is a component of the dipole moment parallel to the chain contour and chain motion is dielectrically active as well. The dynamic measurements were performed utilizing a dielectric spectrometer Alpha Analyser from Novocontrol, in the frequency range 10-2–107 Hz. A disk (hundreds on micrometers) of the material under study was placed within a stainless steel parallel plate capacitor. In the case of the pure polymers, the samples were heated to 120 °C and pressed between two plated electrodes with 30 mm diameter. Fused silica fibers with a diameter of 100 μm were used as spacers. The samples were annealed at 120 °C in vacuum for 24 hours to remove traces of water. In the case of the nanocomposites, the powder was pressed to form disks 12 mm in diameter and 0.3-0.6 mm in thickness. The pellets were annealed at 120 °C in vacuum for 24 hours and adhered between indium foils to improve the electrical contact with the electrodes. Dielectric spectra of the films were measured isothermally in the range -150 to 60 °C. The temperature was controlled via a heated flow of nitrogen gas, by a Quatro Cryosystem. During measurements the samples were kept in a pure nitrogen atmosphere. All samples were cooled to the lowest starting temperature with a cooling rate of 10 °C/min and were thermally equilibrated at successively increasing temperatures before the isothermal data collection.

III. Results and Discussion Structural and thermal characterization Figure 1 shows the XRD and DSC measurements for poly(diethylene succinate), DG-110, in (a) and (b), respectively, as well as those for the poly(ethylene succinate), EG-110, in (c) and (d). DG110 is an amorphous polymer with the XRD measurement (Fig. 1a) showing simply an amorphous halo. DSC (Fig. 1b) confirms the amorphous structure of the polymer, which exhibits a glass


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transition temperature Tg = -26 °C. The respective measurements for poly(ethylene succinate), EG110, reveal a crystalline morphology. The XRD measurements (Fig. 1c) show crystalline peaks at 2θ = 20°, 22.7° and 23.2°, which are in good agreement with the literature for PESu and are assigned to the (021), (121) and (200) planes of its alpha-form, respectively.77

(a)

0

5

EG-110

Intensity (a.u.)

Intensity (a.u.)

DG-110

10

15

o

2 ( )

20

25

(c)

o

20

o

23.2 o

22.7

30

0

5

10

15 o

2 ( )

20

25

1

1

(b)

DG-110

-1

Cp (cal.g . C )

o

-26 C

-1

-1

-2

o

32 C

-1 o

-1 o

0

o

-22 C

0

30

(d)

EG-110

-1

Cp (cal.g . C )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

-50

0

50

o

100

Temperature ( C)

150

-2

o

91 C -50

0

50

o

Temperature ( C)

100

150

Figure 1. X-ray diffraction (a and c) and DSC measurements, expressed as specific heat, Cp, (b and d) for DG-110 (a and b) and EG-110 (c and d). The data in (a) and (c) were measured at room temperature whereas the measurements in (b) and (d) were performed with a heating rate of 10 °C/min.



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Fig. 1d shows the DSC thermogram for EG-110 during the second heating run. It is noted that this specimen was initially heated from room temperature to 150 °C and was subsequently cooled from the melt; both runs were performed with a cooling / heating rate of 10 °C/min. As it will be discussed in Fig.4, a very weak crystallization exotherm at Tc=28 °C together with the glass transition at lower temperatures can be observed during the cooling from the melt. During heating, the glass transition is observed at Tg = -22 °C whereas, above Tg, a main exothermic peak is observed at Tc1 = 32 °C and a second weaker one at Tc2 = 60 °C. Finally, the melting of the material is observed at Tm = 91 °C. This thermal behavior, although not trivial, has been observed before in both isothermal and non-isothermal measurements and its dependence on the cooling rate and on the crystallization temperature was investigated.78 The two most probable explanations for the behavior involve a dual morphology mechanism and the existence of crystals with different thermal stability79 on one hand and a re-organization mechanism of a melting, recrystallization and remelting processes that suggests that the initial crystallization creates crystals of lower degree of perfection or thinner lamellae that melt and recrystallize during heating to produce more perfect crystals.80,81 Alternatively, the possibility that there are two kinds of polymer chains with different mobility in PESu, namely, polymer chains with high molecular mobility and polymer chains with low molecular mobility was proposed; nevertheless, temperature modulated DSC experiments showed that the smaller crystallization exotherm was due to melt-recrystallization of the originally existing unstable crystals formed through previous crystallization.82 A series of nanohybrids was prepared comprised of each of the two polymers and the hydrophilic sodium montmorillonite, Na+-MMT, over a range of concentrations that covered the whole regime between pure polymer and pure inorganic material; this allows the investigation of the structural and thermal properties of the polymers under severe confinement. Figure 2 shows XRD


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measurements of the DG-110 / Na+-MMT (Fig. 2a) and EG-110 / Na+-MMT (Fig. 2b) nanocomposites together with the diffractograms of the pure polymers and of the Na+-MMT. It is noted that all the materials were thermally annealed overnight at 130 °C following the solvent evaporation before the measurements at room temperature. The pure Na+-MMT exhibits a main (001) diffraction peak at 2θ = 8.8º, which corresponds to an interlayer distance of 1.0 nm.

+

+

EG-110/Na -MMT

DG-110 / Na -MMT

o

20

o

o

5.2 =1.7nm

o

66 o

6.4 =1.4nm

45

22.7

o

5.2 =1.7nm

Intensity (a.u.)

DG-110 wt% 100

Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

o

23.2

EG-110 wt%

6.7 =1.3nm

100 93 89 84 67 50

31

30

17

15

0 0

5

10

15

20

25

30

0 0

5

o

2 ( )

10

15

20

25

30

o

2 ( )

Figure 2. X-ray diffractograms of (a) DG-110 (top), Na+-MMT (bottom) and their nanocomposites and (b) EG-110 (top), Na+-MMT (bottom) and their nanocomposites after solvent evaporation and thermal annealing at 130 °C under vacuum for 12h. The measurements were performed at room temperature and are shown shifted along the y-axis for clarity.



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For both polymers, an intercalated structure is obtained for all hybrid compositions evident by the shift of the main diffraction peak of the inorganic material to lower angles. Upon addition of low amount of polymer (e.g., ~15wt%) the peak of the inorganic material has disappeared and a new one at 2θ = 6.4° and 2θ = 6.7° for DG-110 and EG-110 appear corresponding at interlayer distances d001 = 1.4nm and d001 = 1.3nm, respectively. Upon further addition of polymer (i.e., for hybrids with ~30wt% polymer), the existence of a double peak is evident in the diffractograms, with different relative intensity; one is at the same diffraction angle as before whereas the second one appears at lower angles, i.e., at 2θ = 5.2° corresponding to an interlayer distance of d001 = 1.7nm. With further increase of polymer concentration, it is this peak at the lower diffraction angle of 2θ = 5.2° that prevails, always at the same angle independently of the composition. The appearance of those peaks indicates that the polymer chains intercalate between the inorganic layers, forming mono- and bi-layers within the interlayer galleries and at the end a structure with alternating inorganic layers / polymer bi-layers is formed. It is noted that the observed intercalated structures have a significant coherence, which becomes evident by the existence of the third or even the fourth order diffraction peaks. A similar mechanism of chain intercalation in hydrophilic layered silicates has been observed before for poly(ethylene oxide)42 and different generations of hyperbranched polyester polyols.75 Equally important is the observation that, in the case of the EG-110 nanocomposites, the diffraction peaks at high diffraction angles corresponding to the crystalline structure of EG-110 exist only for nanohybrids with very high polymer content (equal or higher than 84wt% in EG-110), indicating that not only the intercalated polymer but the polymer that is close to the stacks of the inorganic particles is completely amorphous and it is only the excess chains that are far away from the inorganic surfaces that are able to crystallize. When the


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crystalline peaks are observed, they are at the same positions with the respective ones of pure EG110 suggesting that the presence of the inorganic material did not alter the crystalline structure.

+

DG-110/Na -MMT

DG-110 wt%

o

100

-26 C

-1

Cp (cal.g . C )

o

-26 C o -29 C

-1 o

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

-1 o

66 45 31

-1

0.5 cal.g . C

-50

0

o

50

100

Temperature ( C) Figure 3. Differential scanning calorimetry thermograms expressed as specific heat, Cp, for DG110 and DG-110 / Na+-MMT nanocomposites with varying polymer concentration; the data are shown for the second heating. The curves have been shifted vertically for clarity. Heating rate: 10 °C/min. The scale bar denotes 0.5 cal.g-1.°C-1. Figure 3 shows the DSC measurements of the nanohybrids composed of DG-110 expressed as specific heat, Cp. As discussed earlier this polymer is amorphous exhibiting a glass transition at Tg = -26 ± 1 °C. It is noted that, for the nanocomposites, the heat capacity has been calculated by taking into account only the mass of the polymer, since the inorganic material is anticipated to contribute only an additive constant to the Cp that does not influence the step of the glass transition. It is evident that the glass transition is observed only in the nanocomposites with a relatively high



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polymer content where the polymer is in excess outside the filled galleries. For composition below ~30wt% in polymer, no transition is observed either because it is too weak to be detected or because it is suppressed when the polymer is confined in the very small dimensions of the interlayer galleries of the inorganic material. Suppression of thermal transitions has been reported in other cases of intercalated polymer nanohybrids of both linear and hyperbranched polymers, as well.70,75,83 When a glass transition is observed, the glass transition temperature is very similar to the one of the pure polymer. The glass transition temperatures of DG-110 and its nanocomposites are reported in Table 2, as well. Table 2. Thermal characteristics of the neat polymers and of their nanocomposites Sample

Tg (°C)

DG-110

-26

66wt%DG-110 + 34wt%Na+-MMT

-26

45wt%DG-110 + 55wt%Na+-MMT

-29

Tc (°C)

Tc1 (°C)

Tc2 (°C)

Tm (°C)

32

60

91

31wt%DG-110 + 69wt%Na+-MMT EG-110

-22

28

93wt%EG-110 + 7wt%Na+-MMT

-19

35

64

94


-19

41

70

96


44

73

97


43

73

97


49



97 98

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+

-1 o

-1

Cp (cal.g . C )

EG-110/Na -MMT

o

EG-110 wt% o 35 C 100 o 93 41 C

28 C

o

44 C o 43o C 49 C -1 o

-1

1 cal.g . C

(a)

o

o

32 C

-22 C

-1 o

(b) -50

-1 o

o

91 C

-1

1 cal.g . C 0

50

89 84 67 50 30

100 93 89 84 67 50

-1

Cp (cal.g . C )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

ο

Temperature ( C)

30 o

97 C 100

Figure 4. Differential scanning calorimetry thermograms expressed as specific heat, Cp, of EG110 and EG-110 / Na+-MMT nanocomposites with varying polymer content during (a) cooling after heating to 150 °C and (b) the second heating runs. The curves have been shifted vertically for clarity. Heating / cooling rates: 10 °C/min. The scale bars denote 1 cal g-1 °C-1.

The thermal behavior of the EG-110 nanocomposites during cooling from the melt with a rate of 10 °C/min and during the subsequent heating is shown in Figure 4 together with the thermograms of the pure EG-110. In contrast to the pure polymer that shows a very weak crystallization exotherm at Tc = 28 °C, the nanocomposites show in general a more pronounced



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transition, the amplitude of which shows a non-monotonic behavior (Fig. 4a); it increases in amplitude and becomes sharper with increasing the concentration of the additive and becomes broader again for the lowest polymer content where it is observed. Irrespectively of the strength of the transition, the crystallization temperature shifts clearly toward higher temperatures with increasing additive concentration and reaches Tc = 49 °C for the 50wt% EG-110 nanohybrid. For the nanocomposite with the lowest polymer content, no crystallization exotherm is observed since all polymer is intercalated and it is impossible to crystallize within the ~1nm galleries of the inorganic material. During heating (Fig. 4b), the glass transition is observed at Tg = -22 °C for the pure polymer; this transition shifts to higher temperatures and becomes extremely weak for the hybrids with polymer concentration higher than ~80wt%. When a glass transition is observed, its temperature is Tg ~ -15 °C. Above the glass transition temperature, the main exothermic peak that is observed for pure EG-110 at Tc1 = 32 °C is not observed in any of the EG-110 nanocomposites. This means that the chains that had crystallized with a cold crystallization in the pure polymer have already crystallized during cooling in the case of the nanocomposites. The second weak exotherm that is observed for pure EG-110 at Tc2 = 60 °C is observed in most of the nanocomposites; nevertheless, the temperature Tc2 shifts significantly towards higher temperatures and reaches the value of Tc2 = 76 °C for the nanocomposite with 50wt% EG-110. The final melting of the polymer, however, is observed at a higher temperature, Tm = 97 °C than that of the bulk EG110; nevertheless, this temperature does not show any dependence on composition. It is noted that, whereas DSC shows crystallization and melting peaks for almost all hybrid compositions, crystalline diffraction peaks are not observed in the XRD measurements for compositions below 67wt% in EG-110 indicating that the structures lack the coherence necessary for a diffraction peak to appear.


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Dynamics The dynamics of the two polymers was investigated in bulk and under confinement close to the inorganic surfaces. Nanohybrids containing 30wt% polymer were utilized to study the effects of confinement since it is estimated that, for the specific compositions, all chains reside within the completely-filled galleries of the inorganic materials. Polymer dynamics was investigated both above and below the glass transition temperature, Tg, of the bulk polymers utilizing Dielectric Relaxation Spectroscopy, DRS; this technique provides the opportunity to study relaxations over a broad range of frequencies and temperatures,. In all cases, the materials were left at 130 °C for 30min, were cooled down to -150 °C with a cooling rate of 10 °C/min and, then, the DRS measurements were performed at the temperatures of interest upon equilibration during heating. Moreover, for the EG-110 specimens, additional measurements were performed following quenching of the sample from 130 °C to -150 °C. Figure 5 shows the imaginary part of the complex permittivity, ", for the bulk polymers with Fig. 5a and Fig. 5c showing the spectra of DG-110 and Fig. 5b and Fig. 5d the respective ones of EG-110. The frequency dependencies of the spectra reveal multiple relaxation processes over the investigated temperature range. For both polymers, at low temperatures (Figs. 5a and 5b) a very broad peak, covering almost seven orders of magnitude in frequency is observed, which shifts to higher frequencies as temperature increases; the breadths of the peaks indicate multiple relaxation processes. These sub-Tg relaxations are related to the motion of hydroxyl groups and/or the orientation fluctuations of the ester groups, which are traditionally referred to as - and processes, respectively.75,84,85 Both polymers exhibit this broad peak at temperatures below the glass transition temperature; nevertheless, differences are observed. For DG-110, the two sub-Tg processes can be identified in the raw data before any analysis at certain temperatures (e.g., at -50


Biomacromolecules

22

°C), whereas a similar separation is more difficult for EG-110 indicating that either the relaxation times of the processes are closer or that the slower one has a lower amplitude.

DG-110

o

EG-110

o

T= -140 C to -50 C

o

o

T= -140 C to -35 C

-1

10

-1

''

''

10

-2

10

-2

10

(b)

(a) 1

1

10

10

-1

-1

o

o

o

o

T= -30 C to -0 C

T= -40 C to 10 C 0

0

10

''

10

''

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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-1

-1

10

10

(c)

(d) -1

10

0

10

1

10

2

10

3

10

4

10

Frequency (Hz)

5

10

6

10

-1

10

0

10

1

10

2

10

3

10

4

10

Frequency (Hz)

5

10

6

10

Figure 5. Imaginary part of the dielectric permittivity, ", as a function of frequency for DG-110 (a, c) and EG-110 (b, d) neat polymers at temperatures of (a) -140 (■), -120 (●), -100 (▲), -80 (▼), -60 (♦) and -50 (◄) °C, (b) -140 (■), -120 (●), -100 (▲), -80 (▼), -60 (♦), -50 (◄), -40 (►) and -35 (●) °C, (c) -40 (■), -36 (●), -30 (▲), -26 (▼), -20 (♦), -10 (◄), 0 (►) and 10 (●) °C and (d) -30 (■), -28 (●), -24 (▲), -20 (▼), -16 (♦), -11 (◄), -8 (►), -5 (●) and 0 (●) °C.

At higher temperatures, near and above the calorimetric Tg, another process is observed as can be seen in Fig. 5c and Fig. 5d; this relaxation is identified with the segmental dynamics (-process)


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23

associated to the glass transition. At even higher temperatures, the observation of any other relaxation processes is obscured by the presence of DC conductivity, which is manifested by the -1 dependence of the " data. 0

10

o

EG-110 EG-110 quenched

''

T = -90 C

-1

10

(a)

-2

10

o

T = -20 C

EG-110 EG-110 quenched

1

10

''

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

0

10

(b)

-1

10

-2

10

-1

10

0

10

1

10

2

10

3

10

4

10

Frequency (Hz)

5

10

6

10

Figure 6. Frequency dependence of the imaginary part of the dielectric permittivity, ", for EG110 cooled with a 10 °C/min rate (■) or rapidly quenched (●) for temperatures below (a) and above (b) the glass transition temperature of the polymer.

To investigate the effect of the cooling rate on the measured dynamics, dielectric measurements were performed for the bulk EG-110 polymer following quenching from the melt. Figure 6 shows



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a comparison of the imaginary part of the complex permittivity, ", for the quenched sample and the one, which was cooled down with a 10 °C/min rate. The comparison is shown for two different temperatures, one below (Fig. 6a) and one just above (Fig. 6b) the calorimetric glass transition temperature of the polymer. As shown in the DSC measurement of Figure 4, this material does not show significant crystallization during cooling with a rate of 10 °C/min; thus, it is anticipated that the two different ways of cooling will not change the morphology of the polymer to a large degree and that they may provide information on the origin of the observed sub-Tg relaxation processes. It is clear from Fig. 6a that, for both cases, two relaxation processes can be observed that have qualitative the same relaxation times. Nevertheless, their dielectric strengths are apparently different; the fast relaxation process has a higher amplitude for the slowly cooled sample whereas it is the slow one which becomes more intense when the polymer is quenched. This observation may indicate that the two relaxation processes are related with a single relaxation motion, which can be partly free and partly restricted; the different cooling rates can affect only the relative populations of the fast and slowly moving groups. Above Tg, the segmental mode can be observed in Fig. 6b, possessing similar amplitude but different relaxation times, being faster for the quenched polymer. This can be attributed to the completely unrestricted motion of the polymer segments when they are quenched compared to the restricted motion of part of them that are tied to the crystallites despite the low degree of crystallinity. Similarly, faster segmental dynamics in initially amorphous poly(l-lactic acid) was observed before compared to the respective semi-crystalline (annealed) ones.86,87 DRS measurements were performed for two nanocomposites with compositions where one can anticipate that all polymer chains are intercalated within the completely filled galleries of the inorganic material. Figure 7 shows the imaginary part of the complex permittivity, ", as a function


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25

of frequency for the two nanocomposites with 30wt% DG-110 or 30wt% EG-110 over a broad frequency range and at similar temperatures with the respective ones of the bulk polymers.

+

DG-110 / Na -MMT -1

o

o

10

o

''

''

T= -140 C to -40 C

o

T= -140 C to -30 C

-1

10

-2

10

-2

10

(a)

+

EG-110 / Na -MMT

(b)

1

1

10

10

-1

-1 0

10

o

0

10

o

o

-1

o

T= -20 C to 0 C

''

T= -40 C to -10 C

''

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

-1

10

10

(c)

(d) -1

10

0

10

1

10

2

10

3

10

4

10

Frequency (Hz)

5

10

6

10

-1

10

0

10

1

10

2

10

3

10

4

10

Frequency (Hz)

5

10

6

10

Figure 7. Frequency dependence of the imaginary part of the dielectric permittivity, ", for 30wt% DG-110 (a, c) and 30wt% EG-110 (b, d) nanocomposites at temperatures of (a) -140 (■), -120 (●), -100 (▲), -80 (▼), -60 (♦), -50 (◄), -46 (►) and -40 (●) °C, (b) -140 (■), -120 (●), -100 (▲), 80 (▼), -70 (♦), -60 (◄), -50 (►), -45 (●) and -30 (●) °C, (c) -40 (■), -36 (●), -32 (▲), -26 (▼), -23 (♦), -20 (◄), -15 (►) and -10 (●) °C and (d) -20 (■), -18 (●), -16 (▲), -14 (▼), -11 (♦), -5 (◄), -2 (►) and 0 (●) °C. Certain differences can be observed in the spectra of the hybrids compared with the ones of the neat DG-110 and EG-110 polymers both below and above the bulk Tg. At low temperatures,



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multiple relaxation processes appear in the spectra similarly to the pure polymers; however, the spectra exhibit different shape and possess lower dielectric strength. At higher temperatures, above the glass transition temperatures of the neat polymers (it is noted that for the specific nanocomposites no glass transition could be observed by DSC), a relaxation process is observed at temperatures similar with the segmental relaxation of the bulk polymer; this process is obscured by another slower process. In all cases, analysis of the spectra is necessary to quantitatively compare the findings. Figure 8 illustrates the quantitative analysis of the dielectric data at two different temperatures for the two polymers and their nanocomposites. The ε*(ω) data were analyzed utilizing the empirical Havriliak-Negami (HN) functions88 𝜀∗ 𝜔

𝜀

∆𝜀 ⁄ 1

𝑖𝜔𝜏

(1)

where HN is the characteristic relaxation time,  = 0' – ' is the relaxation strength of the process and ,  (0 < ,  ≤ 1) describe the symmetric and asymmetric broadening of the distribution of relaxation times. At high temperatures, a term with -1 dependence is necessary to fit the data at low frequencies due to the presence of ionic conductivity. In the bulk, for the fastest process the  parameter is fixed to 1 whereas the  (0.2−0.3) and  (0.5−1) parameters increase with increasing temperature; the values for  (0.2−0.3) signify a broad process. The slower sub-Tg process for DG-110 possesses lower relaxation strength =0.1) and appears to be narrower (=0.6-0.8) as compared to the faster process. For EG-110, this process is quite broad with =0.3-0.45 and  ~0.4. In both cases, the slower sub-Tg process is asymmetric in agreement with the identified -process observed in hyperbranched polyesters.75 At higher temperatures, for the bulk polymers the segmental alpha-process appears; this process for both


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polymers is asymmetric, whereas the  parameter is determined around 0.5−0.8 and  is decreasing with increasing temperature taking values from 7 to 5. Moreover, for EG-110 the relaxation strength of the segmental process decreases around 0C due to cold crystallization and the subsequent decrease of the amorphous polymer content.

DG-110 + DG-110/Na -MMT30%

-1

10

o

-1

10

T = -90 C

''

''

o

EG-110 + EG-110/Na -MMT30%

T = -90 C

-2

10

(a)

-3

10

(b)

-2

10

o

o

T = -11 C

T = -26 C 0

0

10

10

''

''

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

-1

-1

10

10

(c)

-2

10

-2

10

(d)

-2

-1

10

0

10

1

10

2

10

3

10

4

10

5

Frequency (Hz)

10

6

10

7

10

10

-2

10

-1

10

0

10

1

10

2

10

3

10

4

10

Frequency (Hz)

5

10

6

10

7

10

Figure 8. Imaginary part of the dielectric permittivity, ", for DG-110 and for the 30 wt% DG110 / Na+-MMT nanohybrid (a, c) at -90 °C (a) and at -26 °C (c) and for EG-110 and the 30 wt% EG-110 / Na+-MMT nanocomposite (b, d) at -90 °C (b) and at -11 °C (d). The red solid lines are the total fit (together with the conductivity at high temperatures) whereas the dotted and dashed lines are the deconvoluted relaxation processes for the neat polymers and for the nanocomposites, respectively.



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For the nanocomposites of the present investigation, the  parameter of all three processes and for both systems is fixed to 1.0. For the fast sub-Tg process the  parameter varies between 0.250.34 and the dielectric strength  is 0.15-0.3 and 0.3-0.5 for DG-110 and EG-110, respectively, and increases with increasing temperature in both cases. For the slower sub-Tg process a relatively low relaxation strength of ~0.1 is obtained for both nanocomposites, whereas the  parameter is determined to be 0.32-0.45 and 0.4-0.65 for DG-110 and EG-110, respectively. At higher temperatures, the segmental relaxation process appears slower than the respective process of the bulk polymers with shape parameter  around 0.3−0.5 for both nanocomposites whereas the dielectric strength  decreases from 1.4 to 0.35. It is noted that, for the nanocomposites, the influence of conductivity to the spectra is significantly lower but there is a strong interfacial polarization effect, the contribution of which is denoted as a line with slope 0.4-0.6. Figure 9a shows the relaxation map of the neat polymers in an Arrhenius representation. The relaxation times of the three processes that resulted from the analysis of the dielectric data described earlier show different temperature dependencies. The two sub-Tg processes possess Arrhenius temperature dependencies for both DG-110 and EG-110. The activation energies for the fast process are 38±1kJ/mol for DG-110 and 30±1kJ/mol for EG-110, whereas they are 60±2kJ/mol for DG-110 and 59±1kJ/mol for EG-110 for the slow process. The values obtained for the fast process are in agreement with the activation energies obtained for -relaxation processes in other systems, as well;60,62,63,64 however, in most of those works, a single sub-Tg process was identified. Two sub-Tg processes were observed in an investigation of the dynamics of semi-crystalline poly(butylene succinate) and poly(butylene succinate-co-butylene adipate);


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29

however, when the materials were quenched and, thus, amorphous, a single broad sub-Tg process was obtained. The additional slower process for the semi-crystalline polymers had a higher activation energy and was attributed to constrained local motion due to the presence of crystallites.64 In the current work, the slower process possesses higher activation energy; nevertheless, it is observed in the case of the non-crystalline DG-110 as well, suggesting that the constraints are of different origin. A single sub-Tg process identified as -relaxation was observed in the investigation of the dynamics of polylactides. Its shape parameters were =0.3 and =1, whereas the obtained molecular weight dependence of the relaxation times led to the suggestion that it was not of local character.60 In that case, the shape parameters of the segmental process were reported as =0.55 and =0.8. Two sub-Tg processes were observed in the dielectric spectra of three generations of hyperbranched polyester polyols,75 which were identified as the  and  processes, i.e., attributed to the local hydroxyl and carbonyl motions; both processes exhibited Arrhenius temperature dependencies with their activation energies attaining large values because of restrictions imposed on the motion by the network formed due to hydrogen bonds. Under confinement, when the chains were intercalated between the inorganic layers of Na+-MMT, two sub-Tg processes (' and ') were observed as well. However, their activation energies were significantly lower than those of the neat polymers indicating that the two motions are less restricted; this was explained by the formation of fewer hydrogen bonds because of the more flatten conformation that the polymers assume inside the galleries. The activation energy of the -process for the neat hyperbranched polymers is very similar to the one of the current slow sub-Tg process whereas the activation energy of the ' process of the confined chains resembles the one of the fast sub-Tg process of DG-110 and EG-110. The difference between the hyperbranched polyester



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polyols and the linear ones of the current study is that the former possesses many hydroxyl groups in the periphery of the molecules whereas the latter only two end-groups resulting in fewer hydrogen bonds and, thus, to functional groups that move under fewer restrictions or, even, freely. This explanation may also be supported by the DRS measurements of the slowly cooled and quenched EG-110 shown in Fig. 6a, which indicated the exact same relaxation times and only a difference in the relaxation strengths and, thus, in the population of the moving species in the two cases. Thus, the faster process is attributed to free rotation of local (hydroxyl) groups whereas the slower one to restricted motion of similar groups because of the presence of the hydrogen bond network. It is noted that the ester reorientation observed in the investigation of hyperbranched polymers is not observed probably because its relaxation times are very close to the ones of the segmental mode due to the very low molecular weights of DG-110 and EG-110. When the temperature increases above the glass transition temperature, the segmental relaxation of the two polymers is observed. For this process a different behavior is observed, as expected, due to the differences of the calorimetric glass transition of DG-110 and EG-110 with its relaxation times following the Vogel-Fulcher-Tammann (VFT) equation,  = 0 exp[B/(T-T0)]. To minimize the error in the fitting of the data with the VFT equation when the range of the experimental data is limited, the value of 0 = 10-13 is kept fixed for all samples. Fitting of the data resulted in B = 1177±5 K and T0 = 199.0±1.0 K for DG-110 and B = 1286±5 K and T0 = 208.0±3.0 K for EG-110. The Vogel temperature, T0, between the two polymers shows a similar difference as the calorimetric glass transition temperatures, as expected. Moreover, the fragility parameter has similar values D = B/T0 ~ 5.9 and 6.2 for the two polymers DG-110 and EG-110, respectively.


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8

Bulk DG-110 EG-110

6 4 2

fast

0

-log(max/sec)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

-2

slow



(a)

8

Nanocomposites DG-110 EG-110

6 4 2 0 -2 (b) 3

4

5

6

-1

1000/T (K )

7

8

Figure 9. Arrhenius relaxation map for the (a) DG-110 and EG-110 neat polymers and (b) for the nanocomposites with 30 wt% polymer. The sub-Tg processes are shown as filled (fast) and halffilled (slower) symbols and the segmental -relaxation as open symbols. The lines in (b) correspond to the relaxation times of DG-110 and EG-110. Figure 9b shows the relaxation times of the two nanocomposites with 30wt% DG-110 and 30wt% EG-110. The three relaxation processes are also present here as well; nevertheless, their relaxation times exhibit both similarities and differences in comparison with the corresponding dynamics of the neat polymers. The two sub-Tg processes of the confined polymers appear at similar relaxation times with the ones of the bulk DG-110 and EG-110 with both of them following



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Arrhenius temperature dependencies; both processes in the nanohybrids exhibit lower activation energies than those of the neat polymers. The activation energies of the fast process are 28.5±0.5kJ/mol for the DG-110 nanohybrid and 19.6±0.4kJ/mol for the EG-110 one, whereas, for the slow process, they are 36±1kJ/mol for DG-110 and 26.0±0.5kJ/mol for EG-110, indicating that the motions become even less restricted. The obtained relaxation times for the segmental process are slower than the respective ones of the neat polymers following a VFT temperature dependence as well. In previous works of ours, the dynamics of different polymers has been studied under confinement, in intercalated polymer / layered silicate nanocomposites and their relaxation times were compared to the ones of the bulk polymers. Faster, similar or even slower dynamics were observed depending on the system and the specific polymer / surface interactions that may speed up or slow down the relaxation times in conjunction with the geometrical confinement effect that speeds up the dynamics due to the confinement restricting the size of the cooperatively rearranging regions.66,68,69,70,71,75,89 However, a common finding of those studies was that in almost all cases, the temperature dependence of the dynamics changed from a VFT to an Arrhenius one indicating that the confined system appears to be a stronger glass. In the current case, the VFT temperature dependence of the segmental relaxation times is retained and the resulted parameters obtained by the fit are B = 1418±88 K and T0 = 204.0±3.0 K for the DG-110 nanohybrid and B = 1165±20 K and T0 = 222.0±2.0 K for the EG-110 one, resulting in fragility parameters D = B/T0 ~ 7.0 and 5.2 for DG-110 and EG-110, respectively. This may be attributed to the small size of the polymers in this study, which may be comparable to the size of the interlayer distance and, thus, does not result in a significant confinement.


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IV. Conclusions The static and dynamic behavior of two different biobased polyester polyols, one amorphous and one semi-crystalline, in the bulk and in their nanocomposites with a hydrophilic layered montmorillonite were investigated. The nanohybrids were synthesized by solution mixing in water over a broad range of compositions, whereas thermal annealing was applied following solvent evaporation to ensure equilibrium. The morphology of the polymers as well as the structure of the nanohybrids were investigated with X-ray diffraction, which showed that intercalated structure was obtained in all nanocomposites with mono and/or bi-layers of chains within the galleries of montmorillonite. The thermal behavior of all systems was studied with differential scanning calorimetry. DSC reveals that the glass transition, whenever it can be resolved, appears insensitive to the presence of the inorganic material whereas the crystallization of the semi-crystalline polymer depends on the composition of the nanohybrid. Moreover, the dynamics of the polymers both in the bulk and under confinement was probed utilizing dielectric relaxation spectroscopy; the intercalated nanocomposites with ~30wt% polymer were utilized, since it can be calculated that for these hybrids all chains are located within the completely filled galleries of the clay. Both polymers as well as their nanocomposites show two sub-Tg processes and the segmental relaxation. The fast and slow local sub-Tg processes show similar relaxation times and activation energies for the two polymers and for the two nanocomposites with the slower one exhibiting a higher activation energy indicating a constrained motion. The segmental relaxation reveals the effect of the glass transition temperature and becomes slower for both polymers in the nanocomposites retaining, however, the VFT temperature dependence.



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Acknowledgements: This research has been supported by the project “National Research Infrastructure on nanotechnology, advanced materials and micro/nanoelectronics” (MIS 5002772) implemented under the “Action for the Strategic Development on the Research and Technological Sector”, funded by the Operational Programme "Competitiveness, Entrepreneurship and Innovation" (NSRF 2014-2020) and co-financed by Greece and the European Union (European Regional Development Fund). The authors acknowledge the support of COST Action MP1202-HINT (STSM-MP1202-020615-059356).


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References and Notes

(1)

Noordover, B. A. J.; Duchateau, R.; van Benthem, R. A. T. M.; Ming, W.; Koning, C. E.

Enhancing the functionality of biobased polyester coating resins through modification with citric acid. Biomacromolecules 2007, 8, 3860-3870. (2)

Holmberg, A. L.; Reno, K. H.; Wool, R. P.; Epps, T. H. Biobased building blocks for the

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“Structure and Dynamics of Biobased Polyester Nanocomposites”

Krystalenia Androulaki, Kiriaki Chrissopoulou, Daniele Prevosto, Massimiliano Labardi and Spiros H. Anastasiadis 8

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Structure and Dynamics of Biobased Polyester Nanocomposites

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