Morphology, thermal properties and molecular dynamics of syndiotactic polystyrene (s-PS) nanocomposites with aligned graphene oxide and graphene nanosheets

Accepted Manuscript Morphology, thermal properties and molecular dynamics of syndiotactic polystyrene (s-PS) nanocomposites with aligned graphene oxide and graphene nanosheets Stefanos Koutsoumpis, Panagiotis Klonos, Konstantinos N. Raftopoulos, Christine M. Papadakis, Dimitrios Bikiaris, Polycarpos Pissis PII:

S0032-3861(18)30798-5

DOI:

10.1016/j.polymer.2018.08.052

Reference:

JPOL 20859

To appear in:

Polymer

Received Date: 24 July 2018 Revised Date:

21 August 2018

Accepted Date: 23 August 2018

Please cite this article as: Koutsoumpis S, Klonos P, Raftopoulos KN, Papadakis CM, Bikiaris D, Pissis P, Morphology, thermal properties and molecular dynamics of syndiotactic polystyrene (s-PS) nanocomposites with aligned graphene oxide and graphene nanosheets, Polymer (2018), doi: 10.1016/ j.polymer.2018.08.052. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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ACCEPTED MANUSCRIPT

ACCEPTED MANUSCRIPT

Morphology, thermal properties and molecular dynamics of syndiotactic polystyrene (s-PS) nanocomposites with aligned

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graphene oxide and graphene nanosheets Stefanos Koutsoumpis1*, Panagiotis Klonos1, Konstantinos N. Raftopoulos2,3, Christine M. Papadakis2, Dimitrios Bikiaris4, Polycarpos Pissis1 1

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National Technical University of Athens, Heroon Polytechniou 9,15780 Zografou, Greece

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Physik-Department, Fachgebiet Physik weicher Materie, Technische Universität München, James-

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Franck-Str. 1, 85748 Garching, Germany

Department of Chemistry and Technology of Polymers, Cracow University of Technology, ul.

Warszawska 24, 31-155 Kraków, Poland 4

Department of Chemistry, Laboratory of Polymer Chemistry and Technology, Aristotle University of

Thessaloniki, GR-541 24 Thessaloniki, Greece

Abstract

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*corresponding author: [email protected]

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In this work we report morphological, calorimetric and dielectric results for polymer nanocomposites (PNCs) based on syndiotactic polystyrene (s-PS) filled with graphene oxide (GO)

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and graphene by employing melt-mixing. The preparation of the PNCs resulted in aligning of the filler in the polymer matrix as observed by SEM and Wide Angle X-ray Scattering (WAXS). Weak polymer-filler interactions were found, opposite to what has been reported in the literature for atactic polystyrene PNCs. Results by calorimetry (DSC) revealed an increase in crystallization temperature of s-PS upon filler addition, suggesting that the fillers offer additional sites, nuclei, for crystallization. At the same time, filler content has no significant effects on crystalline fraction. The glass transition temperature, Tg, decreases slightly in PNCs, most probably, due to loosened molecular packing of the polymer chains. Taken these re-

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ACCEPTED MANUSCRIPT sults together with those for the calculated rigid amorphous fraction (RAF), which does not vary in a systematic way with the filler loading, we suggest that RAF correlates better with the polymer being bound around the crystals rather than that at the surfaces of the fillers. In dielectric spectroscopy (DRS) next to the main segmental α relaxation of bulk s-PS, related to

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the glass transition, an additional filler-related relaxation (α’) was recorded in the PNCs.

Highlights •

alignment of the graphene based nanoplatelets/nanosheets and polymer crystals in

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the PNCs weak polymer/filler interactions

•

Fillers act as additional crystallization nuclei, RAF correlates with crystals.

•

additional filler-imposed molecular relaxation, α’, in the nanocomposites

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•

Introduction

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Syndiotactic Polystyrene (s-PS), firstly synthesized in 1985, is a semi-crystalline polymer with various attractive properties, such as high melting temperature (270 °C), high crystallization rate, low dielectric constant (~2.5) and good chemical resistance [1,2]. Graphene, on the

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other hand, is a two- dimensional (2-D) single-atom-thick sheet of graphite with carbon– carbon (C–C) bond length of 0.142 nm and shares some impressive properties. It has a

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strength of 130 GPa, Young’s modulus 1 TPa, thermal conductivity ∼5000 W/(mK) and electrical conductivity up to 6000 S/cm [3–5]. It can be produced by mechanical exfoliation, epitaxial growth, Chemical Vapor Deposition (CVD), “unzipping” of CNTs (nano-ribbons) and finally by reduction of graphene derivatives such as graphene oxide (GO) and graphene fluoride. Only the reduction gives bulk quantities of graphene and this is what is used when producing polymer nanocomposites (PNCs) [3–5]. Graphene/GO PNCs have been found to improve barrier properties and anti-corrosion properties [6].

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ACCEPTED MANUSCRIPT For semicrystalline polymers, the 3-phase model suggests that three phases of the polymer matrix co-exist: the bulk/amorphous polymer (mobile amorphous fraction – MAF), the crystalline fraction (CF) and the Rigid Amorphous Fraction (RAF) [7,8]. RAF refers to immobilized polymer at the interfaces with the filler (RAFfiller) and by the crystallites (RAFcrystal). It does not

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contribute to the glass transition in Differential Scanning Calorimetry measurements [9,10] and displays either no [11,12] or slower segmental dynamics in Dielectric Spectroscopy measurements [10,13–17]. RAFcrystal does not contribute to the glass transition or the dy-

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namics at all [7,15,16].

Syndiotactic Polystyrene is known as a semi-crystalline polymer that displays polymorphism

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with four major crystal modifications, namely α, β, γ and δ, and some secondary conformations of the above modifications [18–21]. The crystallization of s-PS with various fillers has been studied. Montmorillonite has been found to decrease slightly the crystalline fraction of s-PS [22]. CNT increase crystalline fraction and enchase β-phase crystallinity [23–25].

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Graphene nanoplatelets serves as an effective nucleating agent for s-PS chains, increasing the crystallization temperature during cooling by 10 °C and slightly the crystalline fraction [26]. About the glass transition, the literature for s-PS is poor, especially for PCNs with gra-

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phene since it is a new material. Wang et al. [26,27] report that Tg is unchanged when gra-

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phene or CNT are introduced in s-PS. The same was reported by Huang et al. [23]. More results can be found in the literature for the more common, atactic Polystyrene (a-PS). In general, strong polymer/filler interactions are observed which increase the Tg and decelerate polymer dynamics. Here we focus on filler based on carbon. Hu et al. [28] in a study of a-PS nanospheres/graphene nanosheets, report an increase of Tg from 101 to 109 °C. This is explained by a restriction of mobility of a-PS due to adhesion on the surface of graphene nanosheets. Thermal stability increases as well. These results were supported by Rissanou and Harmandaris [29] in a simulation work, who report slower dynamics in the a-

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ACCEPTED MANUSCRIPT PS/graphene interface in a distance up to 4 nm, while at short distances the molecules arrange parallel to the graphene sheet. Furthermore, Lee et al. [30] found the Tg and the thermal degradation temperature of the PS with polystyrene-grafted GO PNCs were increased up to 2.8 °C and 23.9 °C, respectively. Han et al. [31] report an increase of Tg when

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GO or graphene are introduced in a-PS via melt mixing. The addition of fullerenes C60 in a-PS slows down the segmental dynamics and increases the Tg as found both by experiment and simulation [32]. Kuilla et al. [33] reviewed the effect of graphene mainly on the electric

interaction of the polymer with the graphene.

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properties of various polymer matrices. For a-PS they report that Tg increases due to strong

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Thus, the results from the literature summarized above show different impact of graphene on Tg for a-PS and s-PS. For s-PS/GO PNCs, on the other hand, no relevant results have been reported, to the best of our knowledge. The present study deals with s-PS melt-mixed with GO and graphene with contents up to 2.5 wt%. Next to scanning electron microscopy (SEM)

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and X-ray scattering (SAXS/WAXS) for morphology, we employ Differential Scanning Calorimetry to follow the crystalline fraction and the glass transition. Dielectric relaxation spectroscopy (DRS) is used to investigate the dielectric α-relaxation, related to the glass transi-

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tion, in order to clarify the nature of the interaction between the polymer and the filler.

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Materials and Experimental Techniques Syndiotactic Polystyrene (s-PS) (Mw = 2.1×105 g/mol) pellets were obtained by Dow Chemicals. The GO was produced through a modified Staudenmaier’s method. In a typical synthesis, 10 g of powdered graphite were added to a mixture of concentrated sulphuric acid and nitric acid while cooling in an ice-water bath. Potassium chlorate powder was added to the mixture in small portions while stirring and cooling. The reactions were quenched after 18 h by pouring the mixture into distilled water, and the oxidation product was washed until a pH 6 was reached. The sample was then dried at room temperature. The graphene nanoparti-

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ACCEPTED MANUSCRIPT cles (GNP) with an average thickness of 10 nm, average platelet diameter of 15 μm and bulk density 0.18-0.25 g/cm3 were supplied by XG Sciences Inc., USA. s-PS PNCs containing GO (0.5, 1.0 and 2.5 wt%) or graphene (0.5, 1.0 and 2.5 wt%) were prepared by melt mixing in a Haake-Buchler Rheomixer (model 600) with roller blades and a

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mixing head with a volumetric capacity of 69 mL. Prior to melt-mixing the nanoparticles were dried by heating in a vacuum oven at 130 °C for 24 h. The two components were physically premixed before being fed in the rheomixer. In order to achieve a better dispersion of

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the nanoparticles in s-PS a RETSCH planetary ball mill (model S100) was used. The s-PS flakes along with the proper amount of nanoparticles were fed into the ‘C’ type stainless steel

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grinding jar with a capacity of 25 mL. Five steel spheres were also added as a grinding medium. The milling was set at 500 rpm for a period of 3 hours. Melt blending was performed after ball milling at 280 °C and 30 rpm for 15 min. For comparison, neat s-PS was treated with the same procedure. During the mixing period the melt temperature and torque were

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continuously recorded. After preparation, each PNC was milled and placed in a desiccator to prevent moisture absorption.

Films of about 500±25 μm thickness were prepared using an Otto Weber Type PW 30 hy-

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draulic press connected to an Omron E5AX Temperature Controller at a temperature of

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280±5oC. The moulds were rapidly cooled by immersing them in water at 20oC. Morphology (filler dispersion) was examined by field emission Scanning Electron Microscopy (SEM) employing a FEI NovaSEM 230 apparatus. The SEM chamber operated at room temperature under high vacuum using a Through Lens Detector (TLD) at a voltage of 3 kV. Regarding sample preparation, pieces from the produced films were immersed in liquid nitrogen and then fractured. SEM images were taken at the fractured surface. Prior to the measurement, a thin layer of gold was deposited on the surface of the sample by sputtering.

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ACCEPTED MANUSCRIPT The morphology studies were complemented by X-ray scattering in a wide range of momentum transfers (q = 0.007 ̶ 2 Å-1) which covers both the Small and Wide Angle region (SAXS/WAXS). The measurements were performed with a Ganesha 300XL SAXS-WAXS system (SAXSLAB ApS, Copenhagen/Denmark) equipped with a GENIX 3D microfocus X-ray

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source and optics together with a three-(scatterless)-slit collimation system. During measurement, the samples were mounted in vacuum at ambient temperature. All images were corrected for cosmic background and parasitic scattering. The measurements were done at

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room temperature, and no thermal protocol was pursued to erase thermal history.

The small angle region of the scattering curves was modeled by a sum of a power law

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=

(1)

which describes scattering from structures larger than the resolution of the instrument, and a Guinier-Porod function to describe to scattering from crystalline lamellae. In the empirical

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model of Guinier-Porod, the scattering intensity is given by [34]:

=

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with

,

1

=

,

3 2

≤

>

!



(2)

(3)

and =

!

6

(4)

ACCEPTED MANUSCRIPT In this model, G is a scaling factor, and d is the Porod exponent, characteristic of the fractal dimension of the particles. The wide-angle region was modelled with a sum of many narrow Gaussian peaks corresponding to the Bragg reflections, and two broad Gaussian-Lorentzian functions describing

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the double amorphous halo of PS. In order to reduce the number of fitting parameters, the model was first applied to the curve of the reference sample (neat sample) which has the

A measure of the crystalline fraction was obtained from: "#$%&' = ()* +

+

-./)01

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where

()* +

()* +

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sharpest peaks, and then the positions of the Bragg reflections were fixed.

is the sum of the areas under the Bragg reflections and

the areas under the two halos.

-./)01

(5)

the sum of

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Differential Scanning Calorimetry (DSC) was employed to study the thermal behavior [35,36]. For DSC measurements an amount of approximately 5 mg of the sample was placed in a Tzero® aluminum pan and measured over the temperature range 40 to 300 °C both during

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heating and cooling with a rate of 10 K/min. The measurements were performed in a Q200

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(TA) calorimeter under high purity Nitrogen as purge gas. The apparatus had been previously calibrated with Indium and sapphires, for temperature and heat capacity, respectively. For the neat s_PS matrix, an additional measurement was performed, following a different thermal protocol in order to determine the heat capacity change during glass transition, ΔCp of the fully amorphous sample: after placing the sample in the aluminum pan, it was heated up to 300 °C and immersed it into Liquid Nitrogen to suppress crystallinity (quenched sample).

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ACCEPTED MANUSCRIPT Thermally stimulated depolarization currents (TSDC) [37] is a special dielectric technique in the temperature domain, characterized by high sensitivity and high resolving power, the latter arising from its low equivalent frequencies (10−4 – 10−2 Hz). TSDC measurements were carried out on samples in the form of films of ~1 mm in thickness. The sample was inserted

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between the finely polished brass plates of a capacitor, placed in a Novocontrol TSDC sample cell and polarized by an electrostatic field Ep (~kV/mm) at polarization temperature Tp = 180 °C for time tp = 5 min. With the field still applied, the sample was cooled down to 20 °C (cool-

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ing rate 10 K/min, under nitrogen flow), sufficiently low to prevent depolarization by thermal energy, then short-circuited and reheated up to 200 °C at a constant heating rate, b = 3

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K/min. Temperature control was achieved by means of a Novocontrol Quatro cryosystem. A discharge current was generated during heating and measured as a function of temperature with a sensitive programmable Keithley 617 electrometer.

Dielectric Relaxation Spectroscopy (DRS) [38] was used to study molecular dynamics using an

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Alpha analyzer with a Quatro cryosystem (Novocontrol) for temperature control. The samples are placed between two brass electrodes, forming a capacitor of known area and thickness equal to the thickness of each sample. Gold electrodes were sputtered on the surface

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for better contact. Two distinct measurements were performed. In one case, the complex dielectric function ε*=ε΄ - iε΄΄ was recorded isothermally in the frequency range from 10-1 to

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106 Hz, at temperatures varying from 100 to 200 °C in steps of 2.5, 5 and 10 K depending on the process followed. In the second case the capacitor was heated with 10 K/min from 90 to 240 °C while the dielectric response was recorded continuously at 100, 10, and 1 kHz (isochronal measurements). The relaxations are evident as peaks in the dielectric loss ε΄΄ spectra, which can be fitted using the Havriliak-Negami (H-N) function[39]:

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ACCEPTED MANUSCRIPT 72 2΄΄ 4 = 5[ ] [1 + 84/4:; - ]=

(6)

where fHN is a characteristic frequency related to the frequency of maximum loss (fmax), Δε is the relaxation strength and a and b are shape parameters. The temperature dependence of

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the time scale of the dielectric response can be followed through the Arrhenius plot (activation diagram, plot of the logarithm of the frequency of the dielectric loss peak against reciprocal temperature) and be further analyzed by fitting appropriate equations. The Arrhenius

E-(+ G F?

(7)

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4.-> ? = 4@ AB C−

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equation is common for describing local dynamics [35]:

where fmax is the frequency of ε’’(f) peak, T the temperature, f0 a pre-exponential constant, Eact the activation energy of the relaxation and k Boltzmann’s constant. The Vogel-TammannFulcher-Hesse equation (VTFH) equation, characteristic of cooperative processes, was used

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to describe the time scale dependence of segmental dynamics on temperature [35,40]: 4.-> ? = 4@

HIJ IKIJ

(8)

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where D is the strength parameter, f0 the pre-exponential frequency factor and T0 the Vogel

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temperature. The apparent activation energy, B, is calculated as B = D×T0 and the fragility index, m, as m =16+590/D similarly to previous work in polymer nanocomposites, for example in ref. [41].

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ACCEPTED MANUSCRIPT Results

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Filler dispersion and orientation

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Figure 1 SEM Images for the samples of s-PS with (a,b) 1% graphene and (c,d) 1% GO at different magnifications. One multilayer sheet of graphene/GO is observed in (a,c) with dimensions of 5-15 μm and the dispersion

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of the sheets in (b,d).

In Figure 1, SEM images for the PNCs with 1% graphene and GO are presented. Both GO and graphene are observed as multilayer sheets with dimensions of 5-15 μm. The different layers are visible through the apparent width of the sheets presented and the ripples forming on their surface (Figure 1a,c) [5,27,28]. These sheets are well dispersed throughout the matrix (Figure 1b,d) and oriented parallel to the surface, as a result of the extrusion and thermopressing of the samples. This shall impose some anisotropy in the samples’ properties de-

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ACCEPTED MANUSCRIPT pending on the axis of measurement, as it will be shown in X-Ray Scattering curves in the following. Both graphene and GO show good dispersion in the matrix at loadings up to 1wt% while for higher content, the filler aggregates.

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Crystallinity and filler orientation by Wide and Small Angle X-Ray Scattering

Figure 2 X-ray Scattering curves for the neat s-PS and the PNCs. The curves are vertically translated for clarity.

The X-ray Scattering Curves in the Wide and Small angle range are presented in Figure 2. . Compared to the neat sample, no additional features are observed in the PNCs in the entire q-range. Furthermore, the good exfoliation of graphene/GO peaks is testified by the lack of additional peaks in the Wide Angle X-Ray Scattering (WAXS) region for the PNCs [3,42], thus we expect only a small number of layers in each GO/graphene sheet. 11

ACCEPTED MANUSCRIPT At q < 0.1 Å-1 (Figure 2) we may observe a knee around 0.03 Å-1 for the neat matrix, presumably due to scattering from crystalline lamellae [43]. This knee appears weaker in the PNCs, because of an enhancement of the intensity at very low momentum transfers (q < 10-2 Å-1), presumably because of scattering from the nanofillers. Interestingly, for the same filler con-

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tent, the scattering curves appear almost identical, irrespective of the type of filler. As mentioned in the Introduction, s-PS is semi-crystalline and displays polymorphism with four major crystal modifications, namely α, β, γ and δ, and some secondary conformations of

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the above modifications [18–21]. The γ and δ forms derive from solvent treatment of s-PS [18,19], such as solvent casting [21,44,45], therefore, we do not expect to detect them in

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the present samples. α form appears after fast cooling or annealing of the glassy matrix while β form appears after cooling from the melt and is more responsive to external nucleating agents [46]. In a work of Wang et al. [26] on PNCs of s-PS/graphene nanoplatelets, the filler was found to promote the β form crystallinity. In general, co-existence of the α and β

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phase is observed [18,25,47]. In this work the samples were rapidly cooled from melt, by

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immersion in cold water.

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Figure 3 WAXS data of the neat s-PS and the 2.5 wt% graphene PNC. The fit of the PNC data is shown as well. The peaks observed are attributed to the crystal lattices of α and β crystal forms as reported in [21]. The scattering curves for all the PNCs were similar to the one presented.

We note that the small peaks observed for the neat s-PS, decrease in size or even vanish in

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the PNCs, while the main complex peak at q = 1.5 Å-1 is significantly enhanced. Whether the size of a peak increases or decreases, it does not correlate with the lattice type (α or β),

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hence we can’t claim that the fillers enhance one or the other. The change of intensity, though, can be attributed to the orientation of the polymer crystallites in the PNCs with respect to the sample surface. It has been observed that in films casted from different solvents, chains have different orientations with respect to the surface and this is reflected in the intensity of the peaks corresponding to the various crystal planes [18,48]. Here, SEM (Figure 1) showed that filler platelets have a preferential orientation, with their planes parallel to the sample surface, i.e. normal to the X-ray beam. We can assume that the platelets act as nucleation agents and the crystals formed around them, are oriented to the platelets. 13

ACCEPTED MANUSCRIPT Then, as a consequence of the orientation of the platelets, the crystals in the composites are also preferentially oriented with respect to the sample surface, whereas they are assumed oriented

in

the

reference

sample.

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randomly

Figure 4 Fitting example for the s-PS + 0.5wt% graphene nanocomposite. The crystalline peaks were fitted using Gaussian terms, the amorphous halos with Gaussian-Lorentzian terms, while a Guinier-Porod term and a power law were used to fit the SAXS region.

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The Crystalline Fraction was calculated as described in detail in the Experimental Section. It increases with 0.5 wt % of the fillers (Table 1), presumably because the fillers act as nucle-

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ating agents [23–26] (to be supported later along with results on crystallization temperature in the Crystallinity section). No significant difference is observed on the effects of the two fillers. Further increase of the content of both fillers does not cause any corresponding increase of crystallinity. As observed also qualitatively, the widths of the peaks tend to increase with filler content indicating a decrease in the crystal size, which is consistent with an extrinsic nucleation [49]. Table 1 Crystalline Fraction CFWAXS, Porod exponent d, power law exponent s, and characteristic size of the lamellae as quantified by the radius of gyration Rg.

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ACCEPTED MANUSCRIPT CFWAXS (wt)

d

s

Rg (Å)

s-PS neat

0.23 ± 0.03

2.86 ± 0.03

2.87 ± 0.09

40.6 ± 0.7

0.5 wt% GO

0.31 ± 0.03

2.48 ± 0.03

2.64 ± 0.05

39.0 ± 1.0

1.0 wt% GO

0.31 ± 0.03

2.57 ± 0.03

2.67 ± 0.08

33.7 ± 2.9

2.5 wt% GO

0.31 ± 0.03

2.66 ± 0.05

2.76 ± 0.04

35.7 ± 1.6

0.5 wt% graphene

0.31 ± 0.03

2.66 ± 0.03

2.78 ± 0.07

37.1 ± 0.8

1.0 wt% graphene

0.30 ± 0.03

2.50 ± 0.13

3.00 ± 0.05

45.1 ± 1.8

2.5 wt% graphene

0.32 ± 0.03

2.74 ± 0.05

2.92 ± 0.03

40.3 ± 1.2

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Sample

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We now turn our attention to the small angle region which was modelled by a power law

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and a Guinier-Porod term as described in the Experimental section (Figure 4). The Porod exponent d is in the range 2.7 to 2.9 without systematic variation with filler content (Table 1). This value, being less than 3, is characteristic of a mass fractal formed by the crystalline lamellae [34]. Their size, as quantified by the radius of gyration, does not show any systematic variation with filler content or type either. The exponent s of the power law is the Porod ex-

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ponent of structures larger than the resolution of the instrument. Again, it has values smaller than 3 and thus corresponds to a mass fractal, and it does not vary systematically with filler content. It should be noted though that the contribution of the power law to the total

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intensity is clearly enhanced with increasing filler content. This could mean that part of it corresponds to scattering from the filler platelets in addition to the scattering from polymer

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crystals.

Crystallinity and glass transition study by DSC For DSC measurements, all samples are heated up to 300 °C in order to erase thermal history. During cooling at 10 K/min (upper panel in Figure 5) the neat s-PS crystallizes at ca. 243 °C while having a second smaller peak at 240 °C, whereas, the PNCs crystallize at ca. 246 °C in a single peak. We calculate the crystalline fraction, CF, using the formula:

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ACCEPTED MANUSCRIPT "# =

7L( 1 − MN × 7P

@@%

× 100%

(9)

where Xf is the filler fraction, ΔHc is the enthalpy of the transition calculated from the area of the peak and ΔH100% = 137 J/g [46] the enthalpy of the 100% crystalline sample. CF increases

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slightly in the PNCs and results are presented in Table 2. The CF values as derived from DSC and X-ray Scattering are similar for the neat sample. The different values for the PNCs between the two techniques (Tables 1 & 2) is of no concern since the samples had different

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thermal history. During heating (lower panel in Figure 5) we observe two melting peaks. The peak at higher temperature has Tm ≈ 270 °C while the lower peak has Tm in the range 262-

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265 °C.

Figure 5 DSC thermograms during cooling (upper panel) and heating (lower panel) for all the samples. The curves have been translated vertically for better presentation.

Table 2 Results from DSC and DRS. On DSC, we report the glass transition temperature Tg, change in heat capacity during glass transition normalized to the polymer fraction ΔCp,n, the Crystalline Fraction, the Mobile

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ACCEPTED MANUSCRIPT Amorphous Fraction and the Rigid Amorphous Fraction. In DRS we note the corresponding Tg for the two relaxations observed above calorimetric Tg and the fragility parameter m.

DSC

DRS aBULK

Tg

ΔCp,n

CF

MAF

RAF

Tg,diel

(oC)

(J/gK)

(wt)

(wt)

(wt)

(oC)

(oC)

(±2)

(±0.02) (±2%) (±25%)

(±25%)

(±2)

(±2)

-

-

-

-

98

133

94

-

104

70

72

66

0.35

94

58

77

40

0.35

94

59

89

38

97

0.29

0.00

1.00

-

PS neat

100

0.14

0.22

0.38

0.40

0.5 wt% GO

98

0.09

0.24

0.24

0.52

1.0 wt% GO

93

0.16

0.23

0.42

2.5 wt% GO

98

0.16

0.23

0.42

0.5 wt% graphene

97

0.16

1.0 wt% graphene

98

0.16

2.5 wt% graphene

96

0.14

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PS neat (quenched)

m

Tg,diel

m

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Sample

a'

0.24

0.42

0.34

87

55

81

37

0.23

0.42

0.35

87

47

81

53

0.22

0.38

0.40

94

78

83

82

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At a first glance, the presence of two melting peaks can be correlated to different crystal forms and a quantitative analysis might be possible, however, this is not that straightforward. The reported values for the melting temperature of the α form, range from 273 to 298

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°C, whereas those of the β form, range from 279 to 320 °C. Multiple peaks are often observed [25,46] which makes it difficult for DSC to distinguish between the two phases. In

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such cases, we may employ XRD [21,46] or FTIR [27] to testify the crystal form being present in the samples. C. H. Su et al. [46] in an extensive characterization of bulk-crystallized s-PS report that the melting temperature of the two phases depends strongly on the thickness of the crystalline lamellae. Due to the cooling rates of the preceding cooling runs (i.e. 10 K/min), we believe that what we report in DSC is mainly β form and maybe some α [25,46] . We should note that the crystal form here is expected to be different from the one reported in X-ray Scattering since we followed a different thermal protocol in DSC. We will not comment any more on the crystallinity type since this is beyond the scope of this research. 17

ACCEPTED MANUSCRIPT In Figure 6a, we present DSC thermograms for all the samples in the glass transition region. The glass transition temperature, Tg, decreases for the PNCs, with 1% GO PNC displaying the lowest Tg, namely, 7 K lower than that of the matrix. In Figure 6b, we present the glass transition for a neat s-PS sample that had been previously quenched from melt in order to obtain

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amorphous sample.

Figure 6 (a) DSC thermograms in the glass transition temperature region for all the samples. (b) comparison of DSC thermograms of a quenched sample and the samples crystallized at fixed cooling rate. Heat flow has been normalized to the mass of the sample.

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ACCEPTED MANUSCRIPT From the heat capacity step during the glass transition we can extract information on the Mobile and the Rigid polymer fractions (MAF, RAF) [9,12]. The values of the measured heat capacity step, ΔCp, are normalized to the amorphous fraction of the material [41] 7"0

T1 − MN U 1 − "#

(10)

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7"0,S =

where Xf is the filler fraction and CF is the crystalline fraction. We calculate the MAF by comparing the ΔCp,n of our semicrystalline polymer PNCs to that of an amorphous sample [41].

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For the fully amorphous value, we use the normalized value derived by the sample during quenching (ΔCp,amorhous,s-PS ≈0.29 J/gK). 7"0,S

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V #=

7"0,-./)01/W

,

X'

1 − "#

(11)

RAF, the fraction of the polymer that is immobile in the vicinity of the crystals or the filler,

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can be estimated by employing a 3-phase model [7,8], based on MAF and CF (RAF + MAF + CF = 1) [9,12]. Results for CF and RAF are presented in Figure 7. CF increases slightly in the PNCs, practically not being affected by neither the amount nor the type of the filler. On the

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other hand, RAF changes with filler, taking larger values for the PNCs with GO. Mean values of RAF tend to change with filler loading, moreover they show different trends comparing

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between GO and graphene. However, considering the respective uncertainties (Fig. 7, Table 2), no significant effects of the fillers on RAF are practically observed.

19

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Figure 7 Crystalline Fraction and Rigid Amorphous Fraction versus filler loading for all samples.

Dielectric studies and molecular dynamics

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TSDC thermograms for neat s-PS and PNCs are presented comparatively in Figure 8, in the temperature range between 20 and -190 °C. The depolarization current density is normalized with the applied electric field, so that results for different samples can be com-

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pared to each other not only with respect to the temperature position of a peak (time scale of the corresponding relaxation), but also with respect to the magnitude of a peak (dielectric

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strength of the corresponding relaxation, Δε). In the temperature range from 80 to 110 °C, i.e. in the range of the calorimetric glass

transition (Figure 6 and Table 2), the αBULK relaxation peak is recorded for neat s-PS. We know that the equivalent frequencies of TSDC and DSC measurements are of similar range [10], so the relaxation peak called αBULK in Figure 8 is dielectrically related with cooperative sPS chain motions during glass transition. We will provide additional evidence for that later.

20

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Figure 8 Comparative TSDC heating thermograms for s-PS PNCs and, for comparison, for neat s-PS (matrix). Indicated are the dielectric relaxations αBULK and α’, related to dynamics of bulk-like polymer (glass transition) and to modified polymer dynamics induced from the presence of the filler, respectively. For comparison with

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DSC, indicated is also the temperature region of the recorded glass transition temperatures, Tg.

In the case of PNCs, except for s-PS + 1 wt% GO, complex dispersions appear in the thermograms in the range from 30 to 110 oC, consisting of mainly two peaks, namely the

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αBULK and α’. The temperature position (~92 oC) and strength of αBULK seem unaffected for the different compositions. Regarding α’, the peak maximum of the relaxation varies for the dif-

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ferent PNCs between 45 and 65 oC in Figure 8, while its strength is for all PNCs larger than that of αBULK. Despite that α’ is faster than αBULK (recorded at lower temperatures than αBULK), it should not arise from the local (secondary) polymer dynamics as it is not present in the neat matrix. For the sample s-PS + 1 wt% GO, α’ is exceptionally strong and dominates the TSDC response in the discussed region. It is obvious that α’ relaxation originates from the presence of the fillers, for example, from polymer spatially confined between filler nanosheets [50] or from polymer bound at the surfaces of the nanosheets [51]. We will come back to this point about α’ in the following DRS results. 21

ACCEPTED MANUSCRIPT In DRS, we expect the α relaxation (αBULK) to demonstrate itself as a peak in the ε΄΄ spectra at temperatures higher than Tg [52]. Indeed, for the neat s-PS (black dots in Figure 9a), a peak is observed at 1 kHz at 120 °C and we attribute it to the α relaxation. The overall signal exhibits quite low ε’’ values, which is expected for PS, due to the lack of polar groups

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in the PS macromolecule [1]. In the PNCs, strong internal field effects [53,54] are observed giving rise to higher overall dielectric response as compared to the response of the unfilled s-PS matrix. The signal in the PNCs, absolute values of ε΄΄, increases due to local fields result-

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ing from GO/graphene sheets acting as mini capacitors [54], a behavior we have observed in previous works also [12,53]. The PNCs with 0.5% filler have signal of the same magnitude,

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thus, similar local fields/distribution of filler are expected. The PNCs with 1% and 2.5% filler have even higher dielectric response, as expected for higher filler loading/more local capacitors. The sample with 1% GO has the highest signal (similar to response in the TSDC, Figure 8) which is related to the better distribution of the filler compared to the samples of higher

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loading as we mentioned in the morphology section.

22

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Figure 9 Imaginary part of dielectric permittivity (dielectric loss), ε΄΄, against frequency at the selected temperature 120 °C. (a) shows the raw experimental data (points) for all samples studied. Indicated in (a) are the main relaxations recorded, namely, αBULK and α’ (details in the text). The arrows suggest an increase of the signal

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due to strong internal fields. (b) shows results by analysis, for the two samples indicated on the plot, by means of individual Havriliak–Negami (Eq. 6) components (lines) for each of the recorded relaxations and straight

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lines for conductivity (or substrate).

Not only the signal of the PNCs increases due to strong internal fields, but an extra contribution is also observed. In Figure 9 we observe an extra peak in ε’’ at around 1 Hz, overlapping the signal of the α relaxation peak at 1 kHz. This extra peak is observed also during isochronal measurements, where we record the dielectric response during a temperature ramp, at fixed frequency. In Figure 10 we follow two distinct peaks for the 0.5% graphene PNC. The first peak is the αBULK relaxation. We refer to the second peak as α’, as in TSDC. We should note that in the TSDC measurements, at the equivalent frequency of 10-3 Hz [10], the α’ re23

ACCEPTED MANUSCRIPT laxation is faster than αBULK (appears at lower temperatures), whereas at higher frequencies the relaxation appears to be slower (appears at higher temperatures). Moreover, in Figure 10, the distance between the two peaks, measured in Celsius degrees, seems to increase from 50 °C at 1 kHz to 60 °C at 100 kHz, which suggest different activation energies for the

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two relaxations. α’ relaxation was observed both by TSDC and DRS and its origin shall be ex-

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plored in the discussion.

Figure 10 Imaginary part of dielectric permittivity (dielectric loss), ε΄΄, against temperature at three distinct frequencies, namely 1 kHz, 10 kHz and 100 kHz for the PNC with 0.5wt% graphene. The response was recorded during heating, 10 K/min.

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We performed analysis of the complex ε’’ spectra by fitting of Havriliak Negami, HN, term for

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each relaxation, αBULK and α’, and a conductivity term (actually a straight line) to compensate for the rise of the signal at low frequencies. The fitting procedure has been described explicitly in prior works [55,56] and an example of fitted curves is found in Figure 9b. Results on the time scale and the strength of the relaxations (actually nominal values for the latter because of the internal field) are quantified in the Arrhenius plot (Figure 11). We complement the Arrhenius plot with thermally stimulated depolarization currents (TSDC) peaks at the equivalent frequency of 10-3 Hz [10] and DSC points of Tg at the equivalent frequency of 100s

24

ACCEPTED MANUSCRIPT [35]. From the VTF fitting, we can calculate the Tg,diel, at the equivalent frequency of 100s

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[35,55]. In Table 2 we note values of Tg,diel and fragility m for the two relaxations.

Figure 11 Arrhenius plot for the relaxations above Tg for the neat sample and the PNCs with 2.5% filler. In the inset, we follow the relaxation strength. The lines represent fittings of the VTFΗ equation (eq. 8).

We observe that the two relaxations have different shape and different fragility parameters.

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α’ in general has low fragility (related to cooperativity), as expected for either adsorbed [51] or spatially confined dynamics in the nanometric scale [57]. Obviously, the latter should be true for systems based on the same type of polymer (in bulk, in the form of PNCs [57], of the

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form of polymers adsorbed on an attractive solid surface [51]), similarly to our case here. The different fragility of the two relaxation results also in the crossing of their traces, thus

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we can now better understand the two relaxation peaks spreading further with frequency in Figure 10 and why α’ was faster in TSDC/low frequencies.

Discussion

Filler alignment and crystallinity form SEM and WAXS results suggest that the filler is well dispersed and oriented in the matrix, with the s-PS crystals aligning to the filler sheets. This is not strange for polymers and we cite a recent simulation work of Harmandaris and co-workers [58] where it is shown how a poly-

25

ACCEPTED MANUSCRIPT ethylene chain is folded in a two-dimensional crystal on the graphene surface. The good exfoliation of graphene/GO peaks is testified by the lack of extra peaks for the PNCs in the Xray scattering curves [3]. By thermal treatment we can have either α or β form crystallinity. During synthesis the sam-

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ples were quenched from the melt. When measured by WAXS, the co-existence of the two possible phases was observed with the α crystallinity being more profound, in agreement with the literature [46,47]. For DSC we erased the samples’ thermal history by heating the

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samples above melting temperature and then we cooled them with a steady rate of 10 °C/min for all the samples to have the same thermal treatment. This thermal protocol results

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in β-form crystals [27,46]. The two melting peaks of β-form crystals (Figure 5) have been interpreted as melting of thinner crystals for the peak at lower T and of the thicker crystals for the peak at higher T [25,46]. In the literature on s-PS with CNTs or graphene, the filler has been reported to act as an external nucleating agent, slightly increasing or decreasing the CF

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and promoting β-form crystallinity [23–26]. Nevertheless, results indicate weak polymer/nanofillers interactions. Fillers slightly increase CF and Tc, i.e. the amount and rate of

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polymer crystallization, respectively (Figure 7) [12,36]. Rigid Amorphous Fraction

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From the DSC results we were able to extract information on RAF (Table 2). RAF is known to exist on the filler surface (RAFfiller) or on the crystals (RAFcrystal). An attempt to distinguish between the two (presented in the Supporting Information), resulted in negative values for RAFfiller. This suggests that, possibly, in our case RAF correlates better, if not uniquely, with RAFcrystal, rather than with RAFfiller [59]. The latter has been found true in recent works of polylactide filled with nanofillers of various geometries (GO included) strongly acting as additional crystallization sites [12] and is in agreement with our earlier comment, that no strong

26

ACCEPTED MANUSCRIPT polymer/filler interactions are present, since such would increase RAFfiller and Tg and would lower ΔCp. We note on Figure 7, that PNCs with GO and graphene have similar CF but different RAF (RAFcrystal as we discussed) however, considering only the mean RAF values for RAF. This im-

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plies that we have different semicrystalline morphology in the samples, which has been observed before [42,60]. Samples with more, but smaller crystals, shall have more RAFcrystal due

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to the increased surface of the crystals. Segmental dynamics and origin of α’

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The relaxation of the neat matrix, αBULK, is attributed to dynamics of the amorphous bulk phase, MAF. αBULK is observed for all PNCs, is fitted with the VTFΗ equation (Figure 11) as expected for segmental dynamics in the range of glass transition [35] and gives a decreasing Tg,diel pattern (Table 2), in agreement with DSC. About the origin of α’ relaxation, two are the prevailing scenarios: it should arise either from polymer bound at the surfaces of the

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nanosheets [51] or from polymer spatially confined between filler nanosheets [50]. The scenario of bound polymer on the surfaces of the nanosheets sounds valid and can be

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well supported. The α’ relaxation is slower than αBULK, as it is common for polymer bound at the filler interfaces [13,61]. Schönhals and coworkers [51] recently published a work with

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adsorbed polymer on a solid substrate and reported dynamics qualitatively similar to αBULK and α’ here, as shown in Figure 11. The adsorbed layer exhibited a non-cooperative character (linear time-scale pattern, similarly to α’ here), which was explained in terms of constrains in polymer chains packing [51]. The adsorbed/bound polymer at the surfaces of the nanosheets would suffice here if it were not for some contraindications, such as the steady or slightly decreasing Tg, the nucleation effect of the filler and the RAF mainly correlating with crystals than with the filler. In the literature on polystyrene, Tg has been reported to remain the same when s-PS is filled with CNTs where polymer/filler interactions were weak 27

ACCEPTED MANUSCRIPT [23,26,27] while it increases in PNCs of atactic PS with either CNTs or graphene where polymer/filler interactions are prevailing [28–33]. In general, we don’t expect strong polymer/filler interaction since this would result to increase of Tg and lowering of ΔCp [62]. The decrease of Tg can be understood/rationalized in terms of loosened molecular packing

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of the polymer chains (i.e. of the MAF) confined to nanometer-scale spaces between filler constraints, most probably, also of constraints due to crystals)[57,63] where we expect faster dynamics, lower Tg and Arrhenius like behavior of the segmental dynamics of the confined

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polymer [64]. In our case, s-PS chains confined between the filler sheets shall have lower Tg, increased ΔCp and faster dynamics. No separate Tg could be distinguished for the confined

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polymer but the decreased single Tg (Table 2) could be superposition of Tg from bulk and confined polymer. α' is characterized by lower fragility (related to cooperativity) than αBULK (Table 2), as expected for confined dynamics [64]. This scenario has some contraindications also. The α’ is slower than αBULK, in a wide frequency range, and with an exception of a work

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on PEO/montmorillonite PNCs reported as confinement, most of the times the confined dynamics do not produce crossing of the two relaxations in the Arrhenius plot, as in our case [57,64].

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Thus, it is not clear, as to whether the origins of α’ in the present work can be explained in

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terms of conventional spatial (geometrical) confinement [63], thus cooperative polymer dynamics in space smaller than the cooperativity length in the bulk [63,64] or as adsorbed polymer dynamics.

Finally, one could suggest that the strong α’ relaxation (Figures 8, 9 and 10Figure 9) could be related with electrical charges trapped at interfaces of different conductivity (Maxwell Wagner Sillars, MWS, effect) [65]. Such interfaces can be formed between the filler and the polymer [53] and/or the crystallized and the amorphous polymer and/or the hard and soft polyrurethane phases [66]. Most probably, the situation here is different, judging from the fact 28

ACCEPTED MANUSCRIPT that for the lower temperatures α’ is recorded faster (at higher frequencies) as compared to the αBULK (Figure 11), which is not expected for a relaxation of the MWS type [65].

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Conclusions In this work we studied nanocomposites of syndiotactic polystyrene (s-PS) and graphene oxide (GO) or graphene. The filler dispersion is good and the GO/graphene sheets are well

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exfoliated. To the best of our knowledge, there are no systematic studies on the glass transition or dielectric properties of such systems to be found in the literature. We report serious

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indications for weak polymer/filler interactions, with the filler slightly increasing the CF and the Tg slightly decreasing. The presence of the filler gives rise to an extra segmental relaxation next to that of the polymer in bulk which may arise either by polymer adsorbed at the filler surface or confined between the filler sheets.

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RAF seems to correlate better with RAFcrystal, confirming the low polymer/filler interactions. The filler is aligned in the matrix, a property with technological interest which can be applied to increase the barrier properties or the conductivity of the samples at selected direction. It

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also acts as an external nucleating agent, with crystals growing in the filler surface, as is was shown by X-ray scattering, affecting slightly the CF when the samples have the same thermal

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history.

Quenching of the PNCs by immersion in Liquid Nitrogen did not result in fully amorphous samples (not presented here), this being another indication that filler act as additional crystallization sites. Faster cooling rates for the PNCs to obtain fully amorphous samples (by using smaller quantities of samples or by employing differential fast scanning calorimetry [67]) shall reveal more information on the polymer/filler interactions.

29

ACCEPTED MANUSCRIPT Acknowledgements We gratefully acknowledge Deutscher Akademischer Austauschdienst for the travel support within the program “Hochschulpartnerschaften mit Griechenland”.

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Crystallization and Homogeneous Nucleation Kinetics of Poly ( ε - caprolactone ) ( PCL ) with Di ff erent Molar Masses, Macromolecules. (2012) 3816–3228.

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