Morphology, crystallization and rigid amorphous fraction in PDMS adsorbed onto carbon nanotubes and graphite

Accepted Manuscript Morphology, crystallization and rigid amorphous fraction in PDMS adsorbed onto carbon nanotubes and graphite Panagiotis Klonos, Iryna Y. Sulym, Dariusz Sternik, Pavlos Konstantinou, Olena V. Goncharuk, Anna Deryło–Marczewska, Vladimir M. Gun'ko, Apostolos Kyritsis, Polycarpos Pissis PII:

S0032-3861(18)30143-5

DOI:

10.1016/j.polymer.2018.02.020

Reference:

JPOL 20368

To appear in:

Polymer

Received Date: 9 October 2017 Revised Date:

11 December 2017

Accepted Date: 12 February 2018

Please cite this article as: Klonos P, Sulym IY, Sternik D, Konstantinou P, Goncharuk OV, Deryło– Marczewska A, Gun'ko VM, Kyritsis A, Pissis P, Morphology, crystallization and rigid amorphous fraction in PDMS adsorbed onto carbon nanotubes and graphite, Polymer (2018), doi: 10.1016/ j.polymer.2018.02.020. 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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Morphology, Crystallization and Rigid Amorphous Fraction in PDMS Adsorbed onto Carbon Nanotubes and Graphite Panagiotis Klonosa,*, Iryna Y. Sulymb, Dariusz Sternikc, Pavlos Konstantinoua,

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Olena V. Goncharukb, Anna Deryło–Marczewskac, Vladimir M. Gun’kob, Apostolos Kyritsisa, Polycarpos Pissisa a

Department of Physics, National Technical University of Athens, Zografou Campus, 15780,

Athens, Greece

Chuiko Institute of Surface Chemistry, National Academy of Sciences of Ukraine, 17 General

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b

Naumov Street, Kiev 03164, Ukraine

Faculty of Chemistry, Maria Curie–Skłodowska University, Maria Curie–Skłodowskiej Sq. 3, 20–

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c

031 Lublin, Poland

* Corresponding author, e–mail address: [email protected]

Abstract

Morphology and thermal transitions of low-molecular weight polydimethylsiloxane (PDMS)

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adsorbed onto carbon nanotubes (CNTs) and micrometric colloidal graphite sheets (CG) were studied employing SEM, isothermal nitrogen adsorption–desorption and calorimetry techniques. The CNTs were found to agglomerate forming voids between tubes of a broad range, while

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adsorption of the polymer from a solution results in the expected wrapping of CNTs by PDMS chains and, further, in filling of voids, as documented by SEM and by reduction in specific surface area, SBET. By employing three different thermal protocols in differential scanning calorimetry

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(DSC) measurements, namely, fast cooling, slow cooling and isothermal annealing of crystallization, direct and indirect filler effects on PDMS mobility (glass transition temperature, Tg, mobile amorphous fraction, MAF), crystallization temperature, Tc, and crystalline fraction (CF) were followed. Results were further evaluated in terms of the rigid amorphous fraction, RAF, and the respective contributions of crystals, RAFcrystal, and of filler-polymer interfaces, RAFfiller. CNTs were found to increase the rate and degree of crystallization of PDMS, while RAF seems to correlate with a previously proposed fraction of polymer ordered at the interface with CNTs for the uncrystallized samples and with CF for the semicrystalline samples. CG particles interact at a lesser extent with PDMS (less RAFfiller), while their presence results in suppression of rate and degree of crystallization (i.e. opposite effects than those of CNTs), so that contributions to RAF could be 1

ACCEPTED MANUSCRIPT separated in that case. For both cases of filler, changes in RAF correlate well with those in the SBET value of initial particles. This result supports previous observations in nanocomposites based on PDMS of various low molecular weights adsorbed onto metal oxide nanoparticles with a wide range of SBET. polydimethylsiloxane;

carbon

nanotubes;

crystallization; rigid amorphous fraction; glass transition.

1. Introduction

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nanocomposites;

polymer

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Keywords:

The use of nanosized fillers in polymer nanocomposites (PNCs) [1] offers the great benefit

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that only a small amount of filler content is sufficient to induce tremendous improvement in desired properties [2], in comparison with those of polymer in bulk. This improvement is, in general,

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achieved at much lower filler factors than in conventional composites [2]. Although there is no theory yet to account for this significant difference, it is generally accepted that interfacial effects play a significant role for this [3]. With the term ‘interfacial effects’ we mean changes in structure and organization, degree of interaction, thermal transitions, molecular dynamics and properties of the polymer at the interfaces with the filler nanoparticles (NPs) extending up to a few nanometers into the polymer matrix [4] (and references therein). For semicrystalline polymers, in addition to

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polymer/filler interactions, improvements in mechanical properties (both in neat matrices and in PNCs) may arise also from crystallinity-related effects [5-7], in particular from an increase in the degree of crystallinity [8-14] (and references therein). One way to increase polymer crystallization is by adding of inorganic inclusions (heterogeneities), carbon nanotubes (CNT) being a

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characteristic example for that [15,16]. It has been found that CNTs added to a semicrystalline polymer matrix, next to causing an increase in the degree and rate of crystallization [8,11-15,17-19],

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they sometimes favor additional morphologies (e.g. trans-crystallinity around CNTs next to the crystals of the bulk-like polymer [17,20] or, in general, an ordering of polymer on their surfaces [8]).

Similarly to other fillers in PNCs, a fine dispersion of CNTs in the polymer matrix is critical to successfully achieve the envisaged final properties for applications [15,19,21-25]. It has been demonstrated that CNTs cannot be easily dispersed in a polymer matrix because their aspect ratio is large enough (above 100–1000), thus, CNTs are severely entangled. Additionally, van der Waals forces between individual CNTs are large in number density and, thus, strong, as a result of the large surface area per unit mass of CNTs. Computer simulations [26] have shown that the shape and dimensionality of the CNT determines the strength, range and type of inter-tube van der Waals

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ACCEPTED MANUSCRIPT attractions and polymer induced repulsions [27]. Macroscopically, entanglements between CNTs in the polymer matrix result in significant increase in viscosity [22,26,28], even at very low CNT loading (lower than 3–4 %), so that mixing becomes difficult and breaking of primary CNTs at large extents often occurs during intense mixing. Disentanglement can be achieved by properly modifying the methods of preparation, so that more energy against the total van der Waals force between nanotubes is applied [15,24]. Good dispersion of CNTs has often been achieved by

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hydrodynamic methods, employing compatibilization methods through the modification of CNT surface [14,21,23,29], and via forced wrapping of CNTs by the polymer [30] (and references therein).

Addition of CNTs into polydimethylsiloxane (PDMS), the polymer of interest in this study

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(and in other rubbers, as well), improves the mechanical, electrical, and thermal properties of the resulting composite [19,31,32] (and references therein). PDMS has been found useful from both the

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applications (in relation to CNTs based PNCs) [31,33-37] and the basic materials research points of view [38-44]. Τhe PDMS–CNT interfacial interactions [31] are dominated by the so called CH-π interactions between the methyl groups of the polymer and the π-electron rich surface of the nanotubes [45]. The CH-π interaction makes an important contribution to the understanding of weak intermolecular forces [26,45]. CH-π interactions are weak hydrogen bonds; they are largely due to dispersion forces and partly to charge-transfer and electrostatic forces [31]. The density of filler-

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particle interactions can be correlated to the amount of PDMS chains adsorbed per gram of CNT after solvent extraction. This quantity has being determined via ‘bound rubber’ tests, while it has also been studied by molecular dynamics (MD) simulations [31]. PDMS chains are conformed as helices due to the corresponding rotations around the Si–O bonds [46]. The elongated helical-like

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conformed PDMS [46] is initially brought in van der Waals contact to the CNT and, then, the polymer chain can wrap around the CNT [20,26,27,30,47].

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Niu et al. [28] studied the structure and rheological properties of various carbon-based particle suspensions of PDMS, in particular, CNTs, graphene, carbon black, and hollow carbon spheres suspended in PDMS, and recorded similar strength in individual bonding between CNT/PDMS and graphene/PDMS. In addition, Bokobza et al. [48] studied the interface between carbon based fillers, namely, CNTs, graphene oxide and graphite, and PDMS by Raman spectroscopy and reported poor interfacial adhesion in all the carbon-based composites with the exception of PDMS/CNT, where it is proposed that wrapping of the tubes by the polymer is the dominant mechanism of adhesion. Furthermore, results revealed that interactions between graphene oxide sheets seem more important than those between the filler and the polymer matrix [48].

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ACCEPTED MANUSCRIPT In most studies of PDMS/CNTs PNCs [22,24,32,34,49,50] the polymer matrix was filled with up to 20 wt. % CNTs [50]. In the present work, the CNT fraction is significantly higher, as we have chosen to prepare composites via physical adsorption of PDMS, at 30 to 50 wt. %, from a low concentration hexane/PDMS (99/1) solution onto dried initial CNTs. In our recent studies [42,5156] on materials prepared in a similar way, however, on the basis of metal oxide particles and physically adsorbed PDMS, results have provided quite interesting information with regard to

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evaluation of polymer-filler interactions, the topology and structure of polymer at the interfacial layer and the respective effects of surface characteristics (roughness, textural porosity) of the hosting particles. The main polymer-particle interaction involved there was the strong hydrogen bonding formed between the hydroxyls on the surfaces of particles and the oxygens in the backbone

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of PDMS [57]. Here, the type of polymer polymer-particle interaction is significantly different (CHπ interaction, in addition to the ability of PDMS to wrap the CNTs, as discussed above). The same

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is true for the shape and dimensionality of the particles, namely, nanotubes and graphite platelets (sheets) against initially spherical particles in the previous studies. Thus, it appears challenging and promising at the same time to employ the methodology of our recent studies on PDMS adsorbed on metal oxide NPs to investigate here effects of structure and of polymer-particle interaction on crystallization and glass transition of PDMS physically adsorbed on CNT and CG micro-sheets. We employ scanning electron microscopy (SEM) and isothermal nitrogen adsorption-

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desorption techniques to study the morphology and porosity characteristics of the systems under investigation. Then, we employ differential scanning calorimetry (DSC) in combination with three thermal protocols, targeting on manipulation of polymer crystallization process, to study the thermal transitions of the polymer in the composites (with focus on glass transition and crystallization) and,

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finally, to evaluate the fractions of the various polymer phases (crystalline, CF, mobile amorphous, MAF, rigid amorphous [58,59], RAF, and the RAF due to the polymer/filler interfaces, RAFfiller [59-

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62], and due to crystals, RAFcrystal [59,60,63-65]), by employing widely adopted models [4,59,60,63]. It has been demonstrated in previous studies that loading of PDMS with CNTs results, in general, in enhancement of crystallization (CF) of PDMS [66], nevertheless, to the best of our knowledge, effects on MAF/RAF have not been reported. Finally, we compare with each other results obtained with PDMS adsorbed on CNTs and on CG micro-sheets and discuss effects in terms of dimensionality and of morphology [18,67,68] of the carbon-based particles.

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ACCEPTED MANUSCRIPT 2. Experimental 2.1. Materials Commercial polydimethylsiloxane (PDMS, Kremniypolimer, Ukraine, linear, –CH3 terminated, code name: PDMS-400, molecular weight Wm ≈5700, degree of polymerization dp = 75) was adsorbed onto multi-walled carbon nanotubes (MWCNT, obtained by catalytic chemical vapor deposition (CCVD) [25] using pyrolysis of propylene on complex metal oxides catalysts) and onto

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commercial Colloidal Graphite (CG, Ukrspecmasla, Ukraine, size of primary particles 2.5-15 µm, prepared from natural flake graphite) at the amount of 30, 40, and 50 wt. % PDMS in dry samples. Before adsorption, the samples were dried at 110 °C for 1.5 h. A hexane solution of PDMS at a constant concentration (1 wt. % PDMS) was prepared and different amounts of the same solution

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were added to fixed amounts of dry carbon materials powder. The suspension was mechanically stirred and finally dried at room temperature for 24 h and then at 80 °C for 3 h. Samples at PDMS

2.2. Experimental and analysis methods

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concentration from 30 to 50 wt. % are in the form of powder, while neat PDMS is liquid.

2.2.1. Scanning electron microscopy (SEM)

The morphology was analyzed using field emission Scanning Electron Microscopy employing a QuantaTM 3D FEG (FEI Company, USA) apparatus. SEM images were taken using

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an Everhart–Thornley Detector (ETD) operating at 5–30 kV in voltage.

2.2.2. Textural characterization

The textural characteristics of initial fillers and of filler/PDMS composites were studied

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employing low-temperature (77.4 K) nitrogen adsorption–desorption isotherms technique [69] using an ASAP 2420N (Micromeritics Instrument Corp., USA) adsorption analyzer. Before

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measurements the samples were outgassed at 110 °C for 24 hours in a vacuum chamber. The specific surface area (SBET) was calculated according to the standard BET method [70] (details of calculations in section A of supplementary material).

2.2.3. Differential Scanning Calorimetry (DSC) Thermal properties of the materials were investigated in helium atmosphere in the temperature range from −175 to 40 oC using a TA Q200 series DSC instrument, calibrated with indium (for temperature and enthalpy) and sapphires (for heat capacity). Samples of ~8 mg in mass were closed in standard Tzero aluminum pans (for powders) or Tzero hermetic aluminum pans (for liquids).

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ACCEPTED MANUSCRIPT Apart from the case of very low polymer molecular weight [55,61,71], PDMS is a semicrystalline polymer [72], characterized by high degree of crystallinity, up to ~0.8 depending on cooling conditions [53,62,71]. Crystallization in PDMS based NCs is significantly affected by the filler [40,51,61]. Our main focus here is on effects by the fillers on thermal characteristics of PDMS, in particular, glass transition and crystallization. Based on previous knowledge, glass transition in semicrystalline polymers (temperature and strength) can be affected by both the

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crystallization and the presence of the filler [18,53,62,65,73,74]. For PNCs based on semicrystalline matrices (such as PDMS) is has been demonstrated [51,60,61,63] that addition of filler affects (mainly suppresses) degree and rate of crystallization. Thus, we attempt here to distinguish between direct and indirect (via crystallization) effects on glass transition, by employing three measurement

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routes in DSC. Beginning with amorphous (melt) samples at the temperature of 40 oC (melt state for PDMS) we employ three distinct thermal routes. (route 1 – rapid cooling) The sample is cooled

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down to −175 oC at a high rate, namely, ~90 K/min on average over the overall temperature range and ~40 K/min in the region of crystallization, aiming at minimization of any effects of crystallization on glass transition [53,71,75], the latter being recorded during the subsequent heating at 10 K/min from −175 to 40 oC. (route 2 – standard cooling) The sample is cooled down to −175 o

C at the fixed rate of 10 K/min and subsequently heated at 10 K/min up to 40 oC. (route 3 –

isothermal crystallization annealing) The sample is cooled rapidly to −70 oC, where it is kept for at

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least 1 hour until completion of crystallization. Then, it is cooled down to −175 oC and, subsequently, it is heated up to 40 oC at 10 K/min. Experimental data were analyzed (crystallization and melting peak maxima, enthalpy values,

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glass transition temperature and change in heat capacity etc.) by employing the supporting TA Universal Analysis software. Base line corrections are performed automatically by the apparatus software.

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Regarding the evaluation of DSC results, firstly, using the enthalpy of crystallization normalized to the polymer content (wPDMS), ∆Ηc,n (Eq. (1)), ∆H c , n = ∆H c , DSC / wPDMS

(1)

the degree of crystallinity (crystalline fraction), CF, was calculated by Eq. (2), CF = ∆H c , n / ∆H100%

(2)

where ∆H100% is the enthalpy of fully crystallized PDMS, taken as 37.4 J/g [71]. As far as glass transition is concerned, the characteristic temperature Tg was determined as the midpoint of the heat capacity step during the transition. Following previous work [42,51-53,

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ACCEPTED MANUSCRIPT 55,61], results for the heat capacity step at glass transition, ∆Cp, were normalized to the same fraction of amorphous polymer according to Eq. (3),

∆C p , n =

∆C pDSC

(3)

wPDMS (1 − CF )

where ∆CpDSC is the measured heat capacity step, taken as the distance between the baselines before and after glass transition at Tg.

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In many cases of polymer nanocomposites (PNCs) studied by DSC, the most striking effect of the nanoparticles is a significant reduction of ∆Cp,n in PNCs [40,42,55,59-61,63,74], interpreted in terms of a rigid amorphous fraction, RAF, i.e. a fraction of polymer being immobilized at the interface with the filler, RAFfiller. Because of the semicrystalline nature of PDMS, an additional RAF

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should exist in close proximity to polymer crystals [58], RAFcrystal. The total RAF, RAFtotal, is thought to be completely immobile in DSC, making no contribution to the glass transition [58-60].

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According to Schick and coworkers [60], a ‘3–phase model’ can be applied for the quantitative estimation of the various polymer fractions from DSC results in semicrystalline PNCs [i.e. crystalline fraction (CF), rigid amorphous (RAFtotal) and mobile amorphous (MAF)]. Thus, we may calculate MAF and RAFtotal using Eqs. (4) and (5), respectively. MAF =

∆C p , n

∆C pPDMS ,amorphous

(1 − CF )

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RAFtotal = 1 − CF − MAF

(4) (5)

∆C pPDMS , amorphous in Eq. (4) is the heat capacity change at glass transition for fully amorphous neat PDMS, found equal to 0.35 J/gK for PDMS-400 here (Table 1). Please note that, according to the

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above equations, both mass fractions refer to the mass of the whole polymer in the composite [factor (1–CF) in Eq. (4)]. Also, since the RAF to MAF ratio may change with temperature

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[60,63,74], we refer to the calculated values as representatives for temperatures around Tg. In previous work on semicrystalline PNCs [42,55,60,63] RAFfiller and RAFcrystal have been disentangled under the assumption that the RAFcrystal to CF ratio, Rc, determined for the unfilled polymer matrix, remains the same in the PNCs. We will perform similar calculations here, bearing in mind, however, that semicrystalline morphology may differ significantly for different samples and different measurement routes [73], as, for example, samples by routes 1 and 2 were crystallized non-isothermally. Thus, the assumption is better justified for samples measured in DSC via route 3, i.e. isothermal crystallization at the same temperature for all samples [4] (and references therein). We estimate RAFcrystal by the term Rc·CF in Eq. (6), where Rc is expected different for the different measurement routes.

RAF filler = RAFtotal − RAFcrystal = RAFtotal − RC ⋅ CF 7

(6)

ACCEPTED MANUSCRIPT We should note that, from the methodological point of view, this method for separating RAFcrystal from RAFfiller has been proved successful in the case of PNCs where the presence of the filler (polymer-particle interactions) suppresses the rate and degree of crystallinity (reduced Tc and CF [42,55,56,60,63]). On the contrary, in the case that the filler particles offer additional crystallization sites (act as nucleating agents, similarly to CNT based PNCs studied here), we have recently proposed on the basis of experimental results that attempting to disentangle the two RAFs

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may not be physically correct [18,73].

An alternative method to disentangle the two RAFs has been described [4], by performing DSC measurements on the same PNCs first in the amorphous state (e.g. by quenching from the melt state) and then in the semicrystalline state (e.g. after crystallization annealing). RAFfiller is obtained

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from the first type of measurements. Next, under the assumption that RAFfiller remains unchanged during crystallization, RAFcrystal is obtained from the second type of measurements, RAFcrystal =

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RAFtotal – RAFfiller. Please note that serious assumptions are also here involved (details in [4]). The method has already been employed in PNCs based on polylactide filled with graphene oxide [65], silica [73] and MgAl layered double hydroxides [74]. Finally, RAFcrystal has been studied by alternative DSC routes [64,76] and by other techniques [77].

3. Results and Discussion

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3.1. Morphology and porosity characteristics

The structural characteristics of the filler play a significant role in determining its properties and the properties of the composite material as a whole. The morphology of CNT/PDMS and CG/PDMS systems was examined employing the SEM method (Fig. 1). The SEM images of the

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CNTs and of the synthesized CNT/PDMS composites are shown in Fig. 1(a,c,e,g). As illustrated in Fig. 1a, the initial CNTs exhibit a network structure. The average diameter of the CNTs is

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determined to about 20 nm. CNTs tend to form agglomerates. For example, it can be seen in Fig. 1a that CNTs in bulk are packed in dense bundles consisting of a large number of randomly arranged CNTs. Fig. 1c demonstrates an increase in average diameter of the CNTs to 40-44 nm due to PDMS adsorption on their surface. These results demonstrate interaction between CNTs and uniform distribution of PDMS over the surface of CNTs. We recall that the main PDMS-CNT interaction is the so called CH-π bond between the –CH3 groups of PDMS and the π-electron rich surface of the CNT [45]. The PDMS chains helix [46] is initially brought in van der Waals contact to the CNT and, subsequently, it can wrap around the CNT surface [26,27,30,47]. This procedure of polymer adsorption is reasonable considering the method of preparation, namely, physical adsorption from a hexane solution of PDMS at low concentration. Coming back to our results, the packing density of

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have been observed before by Raman and FTIR spectroscopies [78] and were attributed there to interfacial PDMS. At PDMS concentration of 40 and 50 wt. % it can be seen in Figs. 1e,g that polymer in the composites is located not only on the surface of CNTs, but also in the voids between them (please compare changes in the contrast in SEM images between Figs. 1e,g and Figs. 1a,c),

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i.e. filling the intermediately formed voids (pores). Figs. 1e,g show good dispersion of CNTs in the

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composites, while the arbitrary arrangement of CNTs remains.

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Figure 1. SEM images for initial particles: (a) CNT and (b) CG, and for the CNT/PDMS composites (left side, c – 30% PDMS, e – 40% PDMS, g – 50% PDMS) and CG/PDMS (right side, d – 30% PDMS, f– 40% PDMS, h – 50% PDMS).

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Colloidal graphite (CG) is a dispersed (in size of sheets) modification of treated natural flake graphite [79]. In contrast to coarse-grained graphite, it is characterized by a powdery finely divided flaky structure, according to the SEM image of Fig. 1b. Initial CG consists of carbon sheets with a length/width of 2.5-15 µm and thickness of several tens of nm. Thus, micro-sized sheets of CG are

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characterized by much lower specific surface area (initially ~2 m2/g) as compared to CNTs. Therefore, even with a polymer content of 30 wt. %, PDMS is expected to interact with the entire

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CG surface, this being supported by the results of Fig. 1d. Due to this interaction, the graphite sheets coated by polymer adhere to each other and the specific surface area of the composite becomes negligible. The gradual covering of CG with increasing the amount of polymer results in the observed (Figs. 1f,h) smoothening of the CG surfaces. At PDMS contents equal to 40 and 50 wt. %, the structure of the composite can be considered as a concentrated suspension of graphite in a PDMS matrix.

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Figure 2. (a) Specific surface area, SBET, against polymer loading, wPDMS, for CNT and CG based samples. Absolute

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values of SBET were added in (a) beside the respective points. The inset to (a) shows in more detail results for the CG based samples. (b) Incremental pore (void) size distributions, IPSD, for initial CNT and CNT/PDMS composites,

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calculated from the nitrogen adsorption–desorption isotherms (details in text).

The pore (void) structure of the samples was analyzed based on the low-temperature nitrogen adsorption/desorption technique (raw data in Fig. SM.1 in supplementary material and results by analysis here in Fig. 2). The CG and CG/PDMS systems were found practically nonporous (Table SM.1 in supplementary material). The initial CNTs and CNT/PDMS composites exhibit the isotherms of type IΙ (H3 type of hysteresis loops in Fig. SM.1 in supplementary material)

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according to IUPAC classification [70,80]. Capillary condensation occurs at pressure p/p0 > 0.85 (due to adsorption in broad mesopores and macropores). The incremental pore size distribution, IPSD, functions (Fig. 2b) show that the textural

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characteristics change after the modification with polymer adsorption (SEM images in Fig. 1). The main contribution to the textural porosity of the initial CNT is due to macropores (i.e. for pore radii 25 nm < R < 100 nm) and secondly due to mesopores (1 nm < R < 25 nm, please see also the values

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of Vmacro/meso and Smacro/meso in Table SM.1 in the supplementary material). CNT/PDMS composites are characterized by bimodal porous structure (Fig. 2b). Furthermore, as compared to initial CNTs, in CNT/PDMS there is a significant decrease in mesopores contribution to the total porosity with a simultaneous increase in contribution of macropores for composites at wPDMS = 30 wt. % as compared to the initial CNT. Thus, at content of PDMS equal to 50 wt. % the porosity sharply decreases (Table SM.1 in supplementary material), as PDMS fills the voids between nanotubes. It is noteworthy that the IPSD peaks (Fig. 2b) in the range of mesopores and macropores (20 - 100 nm) retain practically the same position for all studied samples, whereas they differ in their intensity. Results can be discussed also in terms of the absolute values of the specific surface area, SBET, which gradually decrease from 140 m2/g to 41 m2/g with increasing of PDMS loading (Fig.

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2a) in CNT/PDMS composites. In the case of CG/PDMS SBET is initially quite low (~2 m2/g for initial CG) and is eliminated in the composites (Fig. 2a). In our recent studies on systems based on metal oxide particles (silica and titania) in the form of agglomerates and physically adsorbed PDMS [42,51-53,55] and in the case of high intraparticle porosity silica-gel/PDMS [56] the fraction of the interfacial polymer was found almost proportional to SBET [54]. In those studies SBET has been found representative of the textural surface roughness (in the nanometric scale) of agglomerates and

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seemed to depend on the size of the primary particles (8–80 nm [52,53,55]). These textural pores were proposed to form the initial contact areas/points for the first stages of polymer adsorption [42]. In the present work, we will attempt to check this scenario and the correlation between the fraction of interfacial polymer and SBET of initial particles in systems based on significantly different

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particles than spherical oxide particles in the sense of internal structure (carbon-based material here)

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and shape (2D nanotubes and 1D graphite sheets, D referring to the nanoscopic dimension).

3.3. Differential Scanning Calorimetry (DSC) 3.3.1. Direct effects of filler

Figure 3 shows DSC thermograms for the rapidly cooled samples in the overall temperature range of measurement (Fig. 3a) and in the region of the glass transition (Fig. 3b). The glass transition at about –130 oC, cold crystallization in the range from –110 to –90 oC [51,61,71] and

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melting in the range from –60 to –35 oC are followed in the thermograms. In section B and Fig. SM.2 of supplementary material we show the cooling thermograms of all materials at high rates. By cooling at an average rate of ~40 K/min in the region of crystallization, the elimination of crystallization during cooling (called also hot crystallization [81]) was achieved for pure PDMS and

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CG/PDMS composites, while, on the other hand, crystallization took place for CNT/PDMS. The elimination of crystallization during cooling is directly manifested by the absence of crystallization

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peaks, the latter being clearly seen in the respective cooling thermograms of Fig. SM.2 in supplementary material. In Fig. 3a suppressed cold crystallization during heating is observed for CNT/PDMS as compared to the significant cold crystallization recorded with CG/PDMS. This is expected, as crystallization occurred during cooling for the CNT/PDMS samples. This is the first time of recording hot crystallization of PDMS based composites during rapid cooling [53,62,82]. We have shown that by such rapid cooling procedure neat PDMS of higher and lower molecular weight, MW, namely PDMS-1000 (MW ~8k), PDMS-200 (MW ~4k) and PDMS-20 (MW ~2k) do not crystallize [42], thus, we suggest that the increased crystallizability observed here for CNT/PDMS is a direct effect of CNTs on PDMS. This point will be further discussed in the following in relation to results on glass transition.

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Figure 3. (a) Comparative DSC thermograms during heating for all the samples studied (indicated on the plot) that had been previously cooled from the melt at high rates (rapidly cooled). The recorded heat flow (in Watts) has been normalized to the whole sample mass (in grams). The main thermal events are indicated on the plot. (b) Details in the

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glass transition region. The added lines represent the baselines of the thermograms before and after glass transition.

Coming to the glass transition, the overshoots observed for the linear neat PDMS and the corresponding CG/PDMS composites in Figs. 3a,b are related with structural relaxation [42], not further studied here. In previous work [42], we have showed that these overshoots diminish and gradually disappear with increasing molecular weight of PDMS, as well as with addition of

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nanofillers interacting with the polymer [55]. At temperatures higher than the glass transition temperature, Tg, crystallization during heating (cold crystallization) is observed, due to suppression of hot crystallization during cooling [5,6].

With regard to melting at the higher temperatures of Fig. 3a, complex and double melting

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peaks (Tm1,2 in Fig. 3a and Table 1) are typical for systems containing PDMS [53,62,71], including PDMS/CNTs [66], and involve also melting and recrystallization of metastable crystals [71]. The

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presence of the two crystallization peaks can be explained also by a scenario that involves two crystal populations which exhibit distinct melting kinetics, similarly to the scenario proposed recently for poly(ether ether ketone) [83]. We should note that for CG/PDMS systems an additional endothermic peak, in particular a pre-melting peak, is observed at –59 oC (Tm3 in Fig. 3a and Table 1). It has been shown in previous work that recording of complex melting peaks for PDMS depends on the cooling/crystallization method [53,56,62] (similarly to other polymers [84]), this being checked also in the present study (results and further discussion in a following section, 3.3.2.2). Figure 4 shows the polymer loading dependence of selected quantities of interest from measurements on samples cooled rapidly in DSC, all values being listed in Table 1. The glass transition temperature, Tg, in Fig. 4a increases with polymer loading, this increase being very strong in the case of CNT/PDMS and Tg for these samples (varying from –129 to –124 oC) is in all cases 13

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higher than Tg of neat PDMS (–130 oC). Constraints imposed on polymer mobility could be at the origin of this increase of Tg in the PNCs. These constraints do not arise, however, only from the strong polymer/CNT interactions, but mainly from the significant crystallization observed for the CNT/PDMS systems of 40 and 50 wt. % PDMS loading (CF 0.35 and 0.37, Table 1). On the other hand, different results were recorded for CG/PDMS, where for all amorphous PNCs Tg (–131 and – 130 oC, Table 1) is rather similar to that of neat PDMS (–130 oC). This result suggests that no

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significant constraints are imposed by CG on polymer mobility, whereas a small reduction of Tg (in the range of experimental error) may suggest spatial confinement effect on polymer between

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graphite nanosheets [56,65].

Figure 4. Polymer loading dependence of (a) glass transition temperature, Tg, (b) total rigid amorphous fraction, RAFtotal, and (c) mobile amorphous fraction, MAF, for the rapidly cooled samples. The lines were added in the plots to guide the eyes. The inset to (a) shows the respective dependence of the crystalline fraction, CF, which seems to resemble that of the Tg (details in the text). The insets to (b) provide simplified models for the various polymer fractions in (left) amorphous (2−phase model, MAF/RAF) and (right) semicrystalline samples (3−phase model, MAF/RAF/CF) [58,60,73]. In (c) the marked areas 1-3 correspond to MAF observed previously in conventional PNCs (1) polyamide/layered silicates [60], (2) polylactide/silica [73] and (3) crosslinked PDMS/silica [62], where the MAF was found almost constant, independently of the filler loading.

14

ACCEPTED MANUSCRIPT Next to effects on Tg, it is interesting to follow effects on the rigid (RAFtotal) and mobile (MAF) amorphous fractions in Figs. 4b and c, respectively. We recall that values for these quantities were estimated from the negative/additive contributions of RAF/MAF to ∆Cp (Eqs. (3-5)) and that they refer to the whole polymer mass in PNCs [factor (1–CF) in Eq. (4)]. RAFtotal decreases with polymer loading, as expected. This is clear for CG/PDMS, as CF=0 and, thus, RAFtotal equals

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RAFfiller. Interestingly, RAFtotal ~RAFfiller for CG/PDMS decreases almost linearly with wPDMS (Figs. 4b). In the case of CNT/PDMS RAFtotal equals RAFfiller only for the highly amorphous CNT/PDMS(30%) sample, while for larger wPDMS RAFcrystal should additionally contribute to the total RAF. A similar situation has been observed before with both amorphous and semi-crystalline

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PNCs of poly(L-lactic acid) (PLLA) filled with well dispersed nano-inclusions of various types and geometries (CNTs, graphene oxide nanosheets, silica and Ag spherical nanoparticles) [18], all

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acting as nucleating agents. It was proposed in that work that RAFfiller in the amorphous state should be considered not as the conventional ‘bound polymer fraction’, but as the fraction of polymer with an ordered structure at the initial stages of crystallization [18,73,85-87], this being partly supported by X-ray measurements [18,86,87] (and references in [18]). We recall in this context that Coleman et al. [8] reported a correlation between the formation of an ordered polyvinyl alcohol (PVA) layer at the interface with CNTs and the respective observed mechanical reinforcement.

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Thus, it is expected that in our case PDMS chains wrap CNTs [30,31] at low polymer loadings (Scheme 1a, Fig. 1c), and that addition of more PDMS results in ordered polymer structures at the CNT-PDMS interfaces (Scheme 1b) [8,18], this ordering being possibly related with formation of crystals in the vicinity of (or around) CNTs [20]. For CG/PDMS, on the other

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hand, only mere adhesion of PDMS on the surfaces of the CG sheets (Scheme 1b) [48] is expected.

4b.

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The scenario described above gains support from results for RAFtotal (~RAFfiller, in general) in Fig.

15

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Scheme 1. Simplified model for explaining effects on the mobile amorphous (MAF), rigid amorphous (RAF) and crystalline (CF) polymer fractions by the gradual adsorption of PDMS chains onto (left) CNT and (right) CG particles.

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For CG/PDMS MAF in Fig. 4c increases with wPDMS, approaching in a systematic way that of 100% MAF for neat amorphous PDMS. For the two semicrystalline CNT/PDMS composites MAF equals ~0.23 wt (Table 1). It is interesting to compare this value with MAF obtained with other PNCs (marked areas 1-3 in Fig. 4c), namely, ~0.33 in polyamide/layered silicates [60], ~0.30

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in PLLA/silica [73] and ~0.23 in PDMS/silica [62]. MAF in semicrystalline PNCs has often been found to be independent from filler loading, similarly to our case here. Wurm et al. [60], who first

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made this observation, suggested that crystallization in the pure polymer as well as in the PNCs proceeds until mobility in the polymer fraction reaches a certain limit, and that, independently of the cause of the mobility restriction (RAF due to filler or crystals), crystallization stops when MAF reaches a characteristic value. We will further discuss this point in the following in the frame of results obtained from measurements by the isothermal crystallization route.

16

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Table 1 Quantities of interest from DSC measurements: crystallization temperature, Tc, normalized enthalpy of (hot) crystallization, ∆Hc,n, glass transition temperature, Tg, normalized heat capacity change during glass transition, ∆Cp,n, enthalpy of cold crystallization, ∆Hcc, melting temperatures, Tm1,2,3, melting enthalpy, ∆Hm, crystalline

sample

Tc (oC)

∆Hc,n (J/g)

Tg (oC) (±0.5)

∆Cp,n (J/gK)

∆Hcc (J/g)

Tm3 (oC)

Tm1 (oC)

Tm2 (oC)

PDMS-400 neat

-

0

–130

0.35

20

-

–52

–40

28

CNT/PDMS (30%)

–97

0

–129

0.10

1

-

–52

–44

CNT/PDMS (40%)

–95

13

–125

0.12

0

CNT/PDMS (50%)

–94

14

–124

0.13

0

CG/ PDMS (30%)

-

0

–131

0.23

CG/ PDMS (40%)

-

0

–131

0.25

CG/ PDMS (50%)

-

0

–130

0.28

PDMS-400 neat

–90

6

–130

0.30

CNT/PDMS (30%)

–88

7

–131

CNT/PDMS (40%)

–83

20

CNT/PDMS (50%)

–82

20

CG/ PDMS (30%)

–98

10

CG/ PDMS (40%)

–93

CG/ PDMS (50%)

-96

MAF (wt) (±10%)

RAFtotal (wt) (±15%)

RAFcrystal (wt)

RAFfiller (wt)

0.00

1.00

0.00

0.00

0.00

4

0.00

0.29

0.71

0.00

0.71

CF (wt) (±0.02)

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∆Hm (J/g)

–52

–42

10

0.35

0.22

0.43

*

*

-

–52

–41

11

0.37

0.23

0.40

*

*

6

-59

–49

–41

11

0.00

0.67

0.33

0.00

0.33

9

-59

–52

–41

15

0.00

0.71

0.29

0.00

0.29

6

-59

–52

–41

14

0.00

0.80

0.20

0.00

0.20

14

-

–52

–40

27

0.16

0.71

0.13

0.13

0.00

0

-

–49

-

5

0.18

0.19

0.63

0.14

0.49

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-

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cooling at 10 K/min

rapid cooling

RAFcrystal. Details of calculations are given in the text.

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fraction, CF, mobile amorphous polymer fraction, MAF, total rigid amorphous fraction, RAF, rigid amorphous fraction at interfaces, RAFfiller, and around crystals,

-

0.00

0

-

–49

–43

10

0.53

0.00

0.47

0.43

0.03

-

0.00

0

-

–48

–42

11

0.53

0.00

0.47

0.43

0.03

–131

0.18

4

-59

–49

–10

8

0.27

0.38

0.35

0.22

0.14

5

–131

0.29

4

-58

–52

–41

16

0.13

0.71

0.15

0.11

0.04

4

–131

0.22

5

-59

–59

–40

14

0.11

0.57

0.32

0.09

0.24

17

–70

32

-

0.00

0

-

-

–46

33

0.85

0.00

0.15

0.15

0.00

CNT/PDMS (30%)

–70

4

–134

0.17

0

-

-

–45

4

0.10

0.44

0.46

0.02

0.44

CNT/PDMS (40%)

–70

23

-

0.00

0

-

-

–45

11

0.62

0.00

0.38

0.11

0.27

CNT/PDMS (50%)

–70

23

-

0.00

0

-

-

–45

15

0.64

0.00

0.36

0.11

0.25

CG/ PDMS (30%)

–70

12

–130

0.16

0

-52

–44

–40

5

0.33

0.31

0.36

0.06

0.30

CG/ PDMS (40%)

–70

11

–130

0.15

0

-52

–45

–41

4

0.30

0.30

0.40

0.05

0.35

CG/ PDMS (50%)

–70

19

–129

0.13

2

-53

–45

–40

14

0.52

0.18

0.30

0.09

0.21

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PDMS-400 neat

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isothermal annealing at –70 oC

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Note (*): for the marked samples and condition of measurement (fast cooling) the separation of RAFcrystal and RAFfiller from RAFtotal is impossible based only on the

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results of the present work.

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3.3.2. Effects of crystallization We proceed now with the investigation of the indirect effects of polymer-particle interaction, namely, those arising from crystallization. To that aim, we employed two routes of DSC measurements in order to increase (induce) crystallization in our samples, namely, by cooling at the fixed low rate of

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10 K/min (results in Section 3.3.2.1, Figs. 5-6) and, furthermore, by performing isothermal crystallization (annealing) at the fixed temperature of –70 oC (results in Section 3.3.2.2, Figs. 7-8).

3.3.2.1. Non isothermal crystallization

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During cooling of initially amorphous (melted) samples at 10 K/min crystallization was recorded for all samples in the form of the exothermic heat flow peaks of Fig. 5a. Crystallization for CNT/PDMS is in general stronger (CF, Fig. 6a) and faster (Tc, inset to Fig. 6a) as compared to neat

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PDMS and CG/PDMS. In particular, Tc of PDMS (initially –90 oC, Table 1) increases in CNT/PDMS by 2-8 K (inset to Fig. 6a, Table 1), while, on the other hand, it is reduced by 3-8 K in CG/PDMS. Additionally, it seems that crystallization during cooling of CNT/PDMS was almost completed, as during the subsequent heating in Fig. 5b no cold crystallization was observed. The opposite was found true for neat PDMS and CG/PDMS PNCs (Fig. 5b). Therefore, it becomes clear already from the raw data that CNTs offer additional crystallization sites, in agreement with observations for the samples

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cooled rapidly (previous section 3.3.1).

Figure 5. Comparative DSC thermograms for the CNT/PDMS and CG/PDMS composites and, for comparison, for pure PDMS during (a) cooling at 10 K/min and (b) the subsequent heating at 10 K/min. The arrows added in (a) indicate the temperature position and height for selected non-isothermal crystallization peaks.

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As far as glass transition is concerned, neat PDMS after formation of CF ~0.16 wt demonstrates Tg = –130 oC, same as that of the quenched sample (Table 1), but, as expected, a weakly suppressed ∆Cp,n (0.30 J/gK, against 0.35 J/gK for the quenched sample, Table 1). CNT/PDMS(30%) exhibits a very weak glass transition step (∆Cp,n= 0.08 J/gK in Table 1), while no step is recorded for CNT/PDMS

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at 40 and 50 wt. % PDMS, these strong changes accompanying the additional formation of CF ~0.18 wt (Table 1) during cooling. For the CG/PDMS samples, in general, Tg remains unchanged (–131 oC, Table 1) and ∆Cp,n is suppressed (with the exception of CG/PDMS(40%)), comparing with the rapidly cooled samples. The weak effects on glass transition in the case of CG arise from the formation of CF

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~0.11-0.27 wt (Table 1), however, glass transition was not vanished here. For comparable values of CF in PNCs, namely, in CNT/PDMS(30%) and CG/PDMS(30%), suppression in ∆Cp by CF is more

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mobility in CNT/PDMS than in CG/PDMS.

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strong in the case of CNT (Table 1). These results suggest, again, more severe restriction of polymer

Figure 6. Polymer loading dependence of the (a) crystalline fraction, CF, and (b) total rigid amorphous fraction, RAFtotal,

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for the samples cooled at 10 K/min. The lines in (a) were added to guide the eyes. The inset to (a) shows the respective dependence of crystallization temperature, Tc. Results suggest increased additive contribution of CNTs, as compared to CG, to the rate and degree of polymer crystallinity (a) and to RAF (b). The insets to (b) provide simplified 3−phase models for explaining effects in CG/PDMS (scheme on bottom/left, polymer crystals formed not close to the particles) and CNT/PDMS (scheme on top/right, CNTs are partly embedded in the crystals [17,18]). The straight line in (b) corresponds to the ideal wPDMS dependence of RAFfiller assuming excellent dispersion of the fillers and full coverage of their surface by the polymer.

Figure 6b shows the wPDMS dependence of RAFtotal (Eq. (5)) for all samples after non isothermal crystallization (cooling at 10 K/min). RAFtotal is again systematically larger for CNT/PDMS than CG/PDMS, this result being similar to that for the rapidly cooled samples in section 3.3.1. As expected, 20

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for neat PDMS, RAFtotal corresponds to RAFcrystal. We attempted to disentangle RAFcrystal and RAFfiller from RAFtotal via the assumption/method described in Section 2.2.3 (Eq. (6)), results being listed in Table 1. Results for RAFfiller and the respective dependence from wPDMS (Table 1) were found neither realistic nor systematic for both cases of particles, indicating that the basic assumption for employing

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Eq. (6) should not be correct here. It has been suggested that this assumption and, subsequently, the employed method should be more consistent with results obtained by isothermal crystallization [18,56,60,65] at the same temperature for all samples (results in the next section). Therefore, we will not proceed here with further discussion of other values recorded by DSC (e.g. Tg, MAF, Tm) for the

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samples cooled at 10 K/min.

3.3.2.2. Isothermal crystallization

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We draw now attention to results by isothermal annealing at –70 oC for all samples (details in section 2.2.3, route 3). The temperature of –70 oC is sufficiently high so that formation of crystals does not start before the isothermal recording for all samples according to Fig. 5a, whereas, at the same time, samples do not suffer large supercooling (which would result in a large number of crystallization nuclei) [5]. Images for crystals of neat PDMS are rare in the literature. Sundararajan [72] showed by SEM that the PDMS crystals at –70 oC are spherulites of about 100 µm in size, however, for PDMS of

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larger molecular weights than in our case.

Figure 7a illustrates the time evolution of isothermal crystallization at –70 oC in the form of heat flow vs time for the samples indicated on the plot. The PNCs of 40 and 50 wt. % in CNTs exhibit faster crystallization (shorter time period to peak maximum in Fig. 7a), than neat PDMS and, further,

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than the CG/PDMS PNCs. For CNT/PDMS(30%) the time to peak maximum in Fig. 7a is almost equal to that for CG/PDMS(50%). This is another indication that supplements our previous results on CNTs

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favoring crystallization of PDMS against CG. It is interesting, however, that in the case of isothermal annealing the degree of crystallinity, CF, achieved for CNT/PDMS(50%) is 0.64, while for the annealed neat polymer CF is 0.85 (Fig. 7a, Table 1). The result seems strange at first glance; nevertheless, it is reasonable, considering that in the PNCs the fraction of particles is quite large, so that the extent of constraints imposed on unbound polymer (MAF) is significant, this resulting in limiting of the full-scale polymer crystallizability, the latter being demonstrated by pure PDMS here (and in previous works on PDMS [53,62]).

21

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Figure 7. (a) DSC heat flow vs time during isothermal crystallization at –70 oC for the initially amorphous samples indicated on the plot. The added vertical scale bar (heat flow axis) corresponds to 0.02 Watts per gram of PDMS. The

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achieved crystalline fractions, CF, are listed on the plot. The inset to (a) shows the polymer loading, wPDMS, dependence of the calculated Avrami indices, ni [88,89], with i taking values 1 and 2 corresponding to first (i = 1) and second (i = 2) stages of crystallization, respectively (details in the text). (b) DSC heating thermograms of the annealed samples (solid lines) and, for comparison, of the same samples previously cooled at 10 K/min (non-isothermal crystallization, dotted lines).

Isothermal crystallization (annealing) was further analyzed here by the widely employed

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Avrami method [88,89]. According to this mathematical method any significant change in the slope of the expected linear-like ‘log(time)’ dependence of ‘log [ln (1–CF)

–1

]’ (intermediate step of Avrami

analysis, examples in Ref. 73) during evolution of crystallization, suggests an altering in crystallization route and this can be expressed as a transition from an initial crystallization stage to another one

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[88,89]. Such change takes place here for all samples. Results of the Avrami analysis (intermediate steps of analysis not shown) of all data are quantified in terms of Avrami index ni [88,89], where i takes

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the values 1 and 2 corresponding to 1st and 2nd stages of crystallization, respectively, for all samples (results in the inset to Fig. 7a). For the neat PDMS the Avrami index ni of the first stage (n1 in Fig. 7a) is equal to ~1.6, increases to ~2.5 for CNT/PDMS with 40 and 50 wt. % PDMS and decreases to ~1 for CNT/PDMS(30%). For CG/PDMS n1 is ~1 and does not seem to vary significantly with polymer loading. Considering the relatively low supercooling [75] (~20 K at –70 oC, as the main melting phenomena are located at about –50 oC, Table 1), according to the empirical Avrami method [88], these results should represent sporadic nucleation [89] for neat PDMS and CG/PDMS accompanied by 1D2D crystals growing. On the other hand, results for CNT/PDMS, n1 values approaching 3, in combination with the overall results on crystallization for CNT/PDMS from the previous sections, suggest instantaneously nucleated crystals [14,89]. It is interesting that the sample composition 22

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dependence of n1 follows that of CF (Figs. 6a and 8a) for CNT/PDMS, both values screening the nucleation activity of CNTs. For the second stage of crystallization (n2 in Fig. 7a) the Avrami index is equal to ~2 for the matrix, decreases to ~1.5 in CG/PDMS and to ~1-1.5 in CNT/PDMS. The Avrami index n2 has been proposed to correspond to formation of n2-D axialities around the already existing

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nuclei (instantaneously nucleated crystals [89]). Such axialities could be, for example, one (n2~1) dimensional lamellar packing around the CNTs and two dimensional lamellar aggregates in the case of CG/PDMS [89]. Other scenarios for explaining the various effects on ni could be also considered, thus, we will not further discuss results by Avrami analysis. Nevertheless, we should keep in mind the

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significantly different results observed for the two different carbon based fillers, mainly for n1 (inset to

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Fig. 7a).

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Figure 8. Polymer loading dependence of the (a) crystalline fraction, CF, and (b) rigid amorphous fraction at the interfaces with the filler, RAFfiller, for the samples crystallized isothermally at –70 oC. The lines in (a) were added to guide the eyes. The insets to (a) and (b) show the respective dependence of RAFcrystal (=const*CF) and of RAFtotal, respectively. Results for

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amorphous samples obtained by rapid cooling, namely (grey solid triangles and line 1) for CNT/PDMS and (grey solid squares and line 2) for CG/PDMS, have been added in (b) for comparison. The straight solid line (3) in (b) is a linear fitting to the overall data by annealing, while the added arrows mark changes in RAFfiller imposed by crystallization.

From the results of ∆Cp,n (Table 1) we calculated RAFtotal (Eq. (5)) for the annealed samples. Results for RAFtotal are listed in Table 1 and the respective wPDMS dependence is shown in the inset to Fig. 8b. RAFtotal is, on average, slightly larger for CNT/PDMS than for CG/PDMS. It is interesting that RAFtotal in PNCs is quite similar to that for rapidly cooled samples (Fig. 4b), with the exception of CNT/PDMS(30%). We recall that for the latter PNC CF =0 after rapid cooling, while for the other two CNT/PDMS PNCs CF is ~0.36 (Table 1, inset to Fig. 4a). These results suggest that crystallization in 23

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CNT/PDMS imposes a reduction in RAFtotal. In case of fillers that do not act as nucleation agents, and, therefore, crystals grow away from the polymer/filler interface, it is expected that crystallization imposes an increase in RAFtotal via additional contribution of RAFcrystal [53,62,65,73]. Obviously, this is not the case for CNT/PDMS here.

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We have previously proposed [18,73] that in the amorphous state RAFfiller represents the fraction of polymer immobilized around the nanoparticles that may act as nucleating agents (the CNTs here), only at the early stages of crystal formation (line 1 in Fig. 8b). Effects of pre-ordering at the early stages of crystallization have been reported in CNT/based PNCs [8,17] and in neat polymers [85,86] by

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a combination of various techniques. When crystals are formed during non isothermal crystallization, they are small in size (probably with a wide distribution of sizes) and mainly around the CNTs (heterogeneous nucleation process). The formation of small crystals within the bulk-like PDMS

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fraction in the PNCs is also probable to occur. The formation of such crystals results in high RAFcrystal values (0.43 wt in Table 1) for the PNCs and unrealistic low value for RAFfiller (Table 1), since the majority of the interfacial chains is embedded within the PDMS crystals that have been formed around the CNTs. However, in the case of isothermal crystallization during annealing at –70 oC, larger crystals may be formed, with more homogeneous distribution of size and containing larger fraction of bulk-like PDMS chains, leading to high values of CF. In this case the RAFcrystal is small (0.11 wt) whereas the

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RAFfiller increases again reaching values close to those estimated by the rapid cooling experiments (Table 1). In this context it is interesting to note that after isothermal crystallization of both, neat PDMS and CNT/PDMS PNCs, the melting events recorded during subsequent heating consist of a single melting peak (contrary to multiple melting peaks observed after non isothermal crystallization,

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Figs. 5b, 7b). This result indicates that the isothermal crystallization procedure leads to the formation of crystallites that are more homogeneous in size and/or of the same form, as compared to those formed

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during non isothermal crystallization. Thus, from the methodological point of view, our results suggest that in semicrystalline PNCs the analysis of DSC data in terms of MAF, RAFcrystal and RAFfiller should certainly take into consideration the spatial distribution and forms of the crystallites (semicrystalline morphology) that are formed.

24

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Figure 9. Fraction of total rigid amorphous polymer, RAFtotal, against specific surface area, SBET, of the hosting particles, for

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composites based on 40 wt. % PDMS-400 (circles, this work). Results are compared with those from previous works on PDMS physically adsorbed at 40 wt. % onto (crossed triangles) fumed silica [42] and titania [55], molecular weight of PDMS ~2k there, and (crossed squares) silica and titania, MW~8k [54]. The line is a linear fit to all data, namely, of the present and previous works.

We may now compare results for RAFtotal of the present work with those obtained in previous works on PNCs based on physically adsorbed PDMS. Figure 9 shows results for the dependence of

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RAFtotal (by DSC) on the specific surface are, SBET, of initial hosting particles in the form of agglomerates before polymer adsorption, for PNCs of PDMS loading at 40 wt. %. Interestingly, results for RAFtotal in CNT/PDMS (SBET of initial CNTs ~140 m2/g) and in CG/PDMS (SBET of initial CG ~2 m2/g) agree with the trend obtained in previous works, namely, with findings for PDMS of ~2k in

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molecular weight, MW, and PDMS of MW ~8k adsorbed at 40 wt% onto initial silica [42,54] and titania [55] particles, covering together a wide range of specific surface area [54]. This trend has been

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discussed in terms of (a) similarities of the method of physical adsorption of small gas molecules (here nitrogen) with that of flexible polymer chains from a solution (of PDMS in heaxane) onto the same hosting porous mean and (b) textural porosity, intra-particular porosity and surface nanometric roughness, as parameters involved in polymer/particle interfacial phenomena [42,54-56]. Going back to our raw data, we may comment on melting and respective differences recorded for different fillers and methods of crystallization (Figs. 3a,5b,7b, Table 1). The presence of multiple crystallizations has been discussed in terms of melting and recrystallization of metastable crystals [71] or/and the existence of more than one crystal populations exhibiting distinct melting kinetics [83,90]. The most interesting result with respect to melting arises from the qualitative comparison of the 25

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melting area in the DSC thermograms (Figs. 3a,5b,7b) between the annealed samples, where melting seems to consist of one main peak, and the not-annealed samples, where melting consists of more than one main peak along with recrystallization (exothermic events) that accompany strong pre-melting effects (endothermic spikes). By observing the two peaks (Tm1, Tm2) recorded for the not-annealed

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CNT/PDMS PNCs, it seems that the melting peak located at around Tm2 dominates (or is favored) after annealing (Fig. 7b). Similar effects have been demonstrated also in previous work on PDMS [51,53,56], namely, that after annealing of crystallization the melting pattern becomes more simple. In relation to this, Androsch et al. [84] recorded an increasing in Tm of the single melting peak for

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polylactide (PLA) with increasing of the crystallization annealing time period at different annealing temperatures and discussed results in terms of conformationally disordered crystals and changes in the dimension (thickness) of the lamellar crystals. We may comment further on the comparison of results

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of single melting peaks for the samples crystallized at −70 ºC (heating scans in Fig. 7b) with those of two (at least) peaks observed after cooling at 10 K/min (cooling and heating scans in Figs. 5a and 5b, respectively). In the latter the temperature at which cold crystallization takes place is between −110 and −85 oC, i.e. much lower than −70 ºC, consequently, crystals formed there should be smaller and their Tm should be lower than that of crystals formed isothermally at −70 ºC. This is expected, as melting of very small crystals produce double peaks on heating since it takes place in a range of temperatures in

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which the polymer chains can recrystallize [5-7]. Based only on our data here, we cannot conclude to a solid explanation for our recordings on melting. In addition, the origin of the additional melting peak (Tm3) for the CG/PDMS systems (Table 1) remains unclear. Several issues remain to be further clarified, whereas new questions arise from the present

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study. Employing DSC measurement on CNT/PDMS cooled at significantly higher rates than in the present study [91] could produce fully uncrystallized CNT/PDMS(30%,40%) samples, so that the

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wPDMS dependence of RAFfiller of Fig. 4b could be completed. Next, it could be checked, for example by employing wide angle X-ray scattering (WAXS) [49] at low temperatures, as to whether the crystallization form of PDMS in the CNT/PDMS composites is similar to that in neat PDMS or not, by comparing the angular (2θ) positions of the WAXS peaks related to polymer crystals. In addition, Xray scattering could supply further insight into the so called ‘ordering’ of the polymer [8,17] at the surfaces of CNTs, in combination, also, with FTIR and RAMAN spectroscopy, that have been proved previously useful in identifying interfacial PDMS related alternations in the surfaces of CNTs in conventional PDMS/CNTs PNCs [78].

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4. Conclusions Effects of carbon nanotubes (CNT) and colloidal graphite (CG) micro-sheets on crystallization and glass transition of physically adsorbed polydimethylsiloxane (PDMS) of relatively low molecular weight (~6k) were in the center of interest of this study. Moreover, effects of polymer/particles

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interactions on the degree of adsorption of PDMS onto the high specific surface area (SBET ~140 m2/g) CNTs and low SBET (~2 m2/g) CG were studied employing isothermal nitrogen adsorption–desorption, morphology (SEM) and calorimetry (DSC) techniques.

As expected, CNTs were found to favor crystallization, increasing the rate and degree of

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crystallinity, CF, whereas the opposite has been found for CG. Glass transition temperature, Tg, seems to be dominated by the degree of crystallinity, while spatial confinement may contribute to the small reduction of Tg in the case of CG/PDMS. Regarding mobility of the polymer, the most interesting

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results arise from effects on change in heat capacity, ∆Cp, of glass transition. DSC results of ∆Cp were analyzed and discussed in terms of mobile and rigid amorphous fractions, MAF and RAF [60], respectively. The RAF arising from direct particle-polymer interactions, RAFfiller, was found larger for CNT/PDMS as compared to CG/PDMS. This result correlates with the proposed model for PDMS adsorption on CNTs that involves the wrapping [26,31] of CNT by the quite flexible polymer chains initiated by a CH-π interaction [45], against simple interactions proposed for graphite / graphene /

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graphene oxide. Interestingly, the SBET (of initial CNTs and graphite) dependence of RAF was found to be in quite close agreement with a previously demonstrated RAF(SBET) dependence of PDMS of various molecular weights physically adsorbed on metal oxide nanoparticles (silica, titania). Moreover, the results for RAF obtained by employing three distinct thermal protocols in DSC provide additional

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support to previously proposed scenarios for CNT-based nanocomposites and, more recently, for PNCs with various nano-inclusions, where the inclusions act as nucleating agents [18,73]. In particular,

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RAFfiller (the interfacial polymer fraction) for CNT/PDMS PNCs seems, on the one hand, to correlate with an ordering of interfacial polymer on the surfaces of CNTs for the not crystallized samples [8,17], while, on the other hand, for the semicrystalline samples, to be merged within the crystals [18,73] formed around the CNT [17,20], provided that rather small and heterogeneous crystals are formed during the conventional crystallization procedure. From the methodological point of view, this is the first time that DSC results for CNT/PDMS composites were analyzed and discussed in terms of the three-phase model ‘CF-MAF-RAF’ [60]. Additionally, although here PDMS is the minority in such type of polymer composites and the method of preparation (adsorption of polymer from solution onto agglomerates of initial particles) does not 27

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resemble that of conventional nanocomposites, the employed methodology for DSC measurements and analysis provided physically accepted results. It would be interesting to check in future work results obtained here on interfacial polymer (fraction, structure) and interpretations on the basis of results by other techniques, such as FTIR and X-ray scattering, as well as to extend this study to other polymers

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and other filler types, where filler nanoparticles act as nucleating agents. Acknowledgements

The authors would like to express their gratitude to Dr. Yuri I. Sementsov and Dr. Yuliya M. Bolbukh (Chuiko Institute of Surface Chemistry, National Academy of Sciences of Ukraine, Kiev,

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Ukraine) for providing the initial particles and useful discussions and the graphic designer Mr. Dimitrios Klonos (http://dimitrisklonos.blogspot.gr/) for the illustrations.

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Highlights crystallization and rigid amorphous fraction of PDMS adsorbed on CNT and graphite

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CNTs act as nucleating agents, increasing degree and rate of crystallinity

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RAF arises from ordering of PDMS (wrapping) around CNTs and around crystals

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separation of RAF due to filler, RAFfiller, and to crystals, RAFcrystal, under assumptions

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weaker interactions (lower adhesion and RAFfiller) of PDMS with graphite micro-sheets

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