A strategy to achieve enhanced electromagnetic interference shielding at ultra-low concentration of multiwall carbon nanotubes in PαMSAN/PMMA blends in the presence of a random copolymer PS-r-PMMA

Abstract

A unique strategy was adopted to achieve an ultra-low electrical percolation threshold of multiwall carbon nanotubes (MWNTs) (0.25 wt%) in a classical partially miscible blend of poly-α-methylstyrene-co-acrylonitrile and poly(methyl methacrylate) (PαMSAN/PMMA), with a lower critical solution temperature. The polymer blend nanocomposite was prepared by standard melt-mixing followed by annealing above the phase separation temperature. In a two-step mixing protocol, MWNTs were initially melt-mixed with a random PS-r-PMMA copolymer and subsequently diluted with 85/15 PαMSAN/PMMA blends in the next mixing step. Mediated by the PS-r-PMMA, the MWNTs were mostly localized at the interface and bridged the PMMA droplets. This strategy led to enhanced electromagnetic interference (EMI) shielding effectiveness at 0.25 wt% MWNTs through multiple scattering from MWNT-covered droplets, as compared to the blends without the copolymer, which were transparent to electromagnetic radiation.

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Introduction

Blending two polymers often generates multiphasic materials with a unique combination of properties. However, the final properties are strongly contingent on the morphology generated during processing and the specific interactions between the components.^1–5 Most often, mixing two polymers results in phase separation at the macroscale due to the chemical incompatibility arising from the large interfacial tension between the components. However, some polymer pairs are miscible under certain conditions and phase separate either upon heating (LCST, lower critical solution temperature) or cooling (UCST, upper critical solution temperature).^6,7 The bi-phasic microstructures thus generated by blending are in general not at equilibrium. Hence, these structures need to be stabilized in order to realize the unique combination of properties that blending offers. The classical approach to suppress coalescence in polymer blends is the use of a block copolymer which reduces the interfacial tension between the components.^8,9 More recently, nanoparticles localized either at the interface or in a specific phase have also shown promising results in suppressing coalescence in polymer blends.^10–12 The latter strategy offers a unique tool to introduce new properties besides stabilizing bi-phasic structures. Hence, this approach has attracted a great deal of interest in literature.^13,14

Among the different nanoparticles, carbon nanotubes (CNTs) have drawn specific attention owing to their extraordinary combination of properties. Hence, understanding the various aspects of polymer/CNT composites has been the focus of many studies; for a review see.¹⁵ The ability of electronic charge transport in CNTs has made them a potential candidate for developing conducting polymer based composites at a relatively low particle fraction. The latter advantageous characteristic results from their large aspect ratio.

In a recent study, we have demonstrated that the localization of CNTs in bi-phasic blends can be tailored using phase separation as a tool.¹⁶ Whereas the CNTs are dispersed randomly in the homogeneous blends, during phase separation, they migrate to their preferred phase.¹⁷ This strategy for selective localization has led to the development of highly electrically conducting materials at relatively low fractions of CNTs. The percolation concentration can be reduced even further by selective localization of the CNTs at the interface of a co-continuous blend.¹⁸ However, this strategy requires a precise control over the mixing time and intensity.

The use of electronic equipment has increased enormously in recent years. As a result, interference of electromagnetic (EM) radiation is a matter of great concern. Hence, shielding the EM radiations has become crucial in order to avoid interference with nearby electronic devices. In this regard, owing to their high strength to weight ratio and good processability, polymer-based conducting nanocomposites have attracted a great deal of interest as EM shielding materials.¹⁹ However, the ease of processability greatly depends on the concentration of the nanoparticles. Achieving high electrical conductivity at low filler concentration is a key challenge.

In a recent review article, Pawar et al.²⁰ discussed the various factors that decide the localization and dispersion of nanoparticles in polymer composites and bi-phasic blends and the resulting mechanism of shielding EM radiation. It is well understood that in order to absorb microwave radiation, materials should possess nomadic charge carriers as well as electrical and magnetic dipoles. In this context, a large amount of research has been dedicated to investigating the potential of various conductive or dielectrically active materials like carbonaceous nanoparticles, conducting polymers and ceramic nanoparticles as well as particles possessing magnetic dipoles like ferromagnetic metal nanoparticles and ferrite particles as fillers in polymer composites for EMI shielding in the GHz frequency range.²⁰ It is now understood that the mechanism of shielding is contingent upon the state of dispersion of the conducting nanoparticles in the polymer composites. For instance, the incoming EM radiation can be reflected or absorbed depending on whether the filler concentration is below or above the percolation threshold. In this context, polymer blends filled with conducting nanoparticles can be an ideal choice for EM wave absorption as the percolation threshold can be tuned in the blends by selective particle localisation.^18,21–23 For instance, Sundararaj et al.²⁴ observed that at 10 vol% loading of carbon black, the EMI shielding and electrical conductivity of PP/PS (polypropylene/polystyrene) blends were lower than those of the PP composites and higher than those of the PS composites. Selective localization of particles can lead to the development of highly electrically conducting materials at relatively low particle fractions which in turn can result in high EMI shielding.

In addition to enhancing the EMI shielding effectiveness at low filler content by selectively localizing particles in one phase of a polymer blend, it has been shown that addition of a copolymer can further enhance the blend conductivity and shielding effectiveness. The responsible mechanisms are a reduction in domain size due to the compatibilizing action of the copolymer, which facilitates the formation of a conductive path throughout the blend, and an improved dispersion of the MWNTs.²⁵ In the present work, we illustrate how a copolymer can be used to tune the localization of MWNTs by tailoring the mixing strategy and components to allow polymer wrapping of the MWNTs. Thereto, we adopted a unique approach in which pristine MWNTs (MWNTs) were wrapped with a random copolymer (PS-r-PMMA) in the first mixing step and subsequently diluted with PαMSAN/PMMA blends with an aim to drive the MWNTs to the interface. The MWNTs, owing to their large aspect ratio, are also observed to bridge the droplets in the blend with the PS-r-PMMA copolymer. This phenomenon is quite interesting from electrical charge transport point of view.

Materials and methods

Both PαMSAN (Luran KR2556) and PMMA (Lucryl G77) were obtained from BASF. An overview of the characteristics of these polymers and of the phase behavior of the blends has been extensively discussed in ref. 16. The PαMSAN/PMMA blends are known to form an LCST system with a phase separation temperature of approximately 190 °C for the used blend composition. The pristine MWNTs were obtained from Nanocyl (Belgium). PS-r-PMMA was procured from Polysciences. The random co-polymer has a molecular weight of 270 000 g mol⁻¹ and consists of 70 wt% styrene and 30 wt% methyl methacrylate with a glass transition temperature (T_g) of 102 °C. A master batch of PS-r-PMMA and MWNTs in a specific ratio (95 : 5, wt/wt) was prepared using melt mixing at 190 °C and 60 rpm for 20 min. Two different sets of blends; 85/15 PαMSAN/PMMA with 5 wt% PS-r-PMMA/MWNTs and 85/15 PαMSAN/PMMA with 0.25 wt% MWNTs were prepared under the same melt blending conditions. Sample abbreviations are summarized in Table 1. It is important to note that the blends were melt mixed at 190 °C which is the onset of demixing in 85/15 PαMSAN/PMMA blends. However, the extruded strands were hot pressed twice at a lower temperature (160 °C). This temperature is below the cloud point temperature of 85/15 PαMSAN/PMMA blends and the applied compression conditions have been confirmed to result in mono-phasic blends.²⁶ Finally, the blends were annealed at 220 °C for 5 hours to allow phase separation to proceed and the MWNTs and random co-polymer to redistribute within the blend.

Table 1 Sample description with effective nanoparticle concentration

S. No.	Sample description	Abbreviation
1	85/15 (wt%/wt%) PαMSAN/PMMA neat blend without filler	85/15 PαMSAN/PMMA
2	85/15 (wt%/wt%) PαMSAN/PMMA melt mixed with 5 wt% of PS-r-PMMA modified MWNTs (PS-r-PMMA : MWNT = 95 : 5, wt/wt); effective concentration of MWNTs becomes 0.25 wt% in the blends	85/15 PαMSAN/PMMA with 5 wt% PS-r-PMMA modified MWNTs
3	PαMSAN/PMMA 85/15 (wt%/wt%) melt mixed with 0.25 wt% MWNTs	PαMSAN/PMMA with 0.25 wt% MWNTs

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Characterization

Phase separated morphology and the localization of MWNTs in the blends was evaluated using transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM). For this, the blend samples were annealed for 5 h and subsequently quenched to arrest the phase microstructure. For morphological studies, ultrathin sections of 100 nm were sectioned with a diamond knife mounted in an ultramicrotome (LEICA EM UC 6). The ultrathin sections were mounted on a Copper-TEM grid of 150 mesh with holey carbon film. The images of the section were captured by TEM (Philips CM 200 FEG) and SEM (FEI Titan 60-300) in bright field transmission mode at an accelerating voltage of 20 kV.

Conductivity spectroscopy measurements were performed at 220 °C on compression molded samples (16.5 mm diameter and 1.75 mm thickness) in the frequency range of 10⁻² to 10⁷ Hz using a Novocontrol Alpha high resolution dielectric analyzer. The instrument measures the complex impedance of the sample and together with the knowledge of the geometry of the sample, the complex conductivity is calculated. Further, in this work, the real part of the complex conductivity will be used. The samples were placed between two brass plates, separated by a circular Teflon spacer to maintain a fixed sample geometry in the melt state. The same spacer was used for all the measurements in order to eliminate systematic errors in sample dimensions. The temperature was controlled by a Novocontrol Quatro temperature controller, which uses a nitrogen gas flow to control the sample temperature with an accuracy of 0.1 °C.

Viscoelastic properties of the blends were studied using a stress controlled Discovery Hybrid Rheometer (DHR-3, TA Instruments) with parallel plate geometry (25 mm in diameter and 1 mm gap distance). A strain amplitude, which is within the linear viscoelastic region was applied to determine the dynamic storage and loss moduli. Isochronal dynamic temperature ramp measurements were performed for evaluating the phase separation temperature of the blends. All the experiments were performed at a uniform heating rate of 0.5 K min⁻¹ from the miscible one phase to the phase separated regime to detect the onset of phase separation in the blends. Isothermal dynamic time sweeps were done at 220 °C to determine the kinetics of morphology evolution from the behavior of the storage modulus (G′) with time in the phase separated regime.

For EMI measurements, an Anritsu MS4642A Vector Network Analyzer (VNA) with coax set-up (Damaskos M 07T) was employed. The scattering parameters were measured in the X and the Ku band frequency limits. A full two port SOLT calibration of the setup was done prior to sample measurements.

Results and discussion

Blend morphology and localization of PS-r-PMMA modified MWNTs at the interface

In our study, we used 85/15 PαMSAN/PMMA blends which are known to form an LCST system with a phase separation temperature of approximately 190 °C for the used blend composition. A random PS-r-PMMA copolymer was selected since PS forms a UCST blend with SAN and thus, under the melt blending and annealing conditions adopted in this study (190 °C and 220 °C), the PS part of the PS-r-PMMA copolymer is expected to be miscible with the PαMSAN phase. Hence, the random PS-r-PMMA copolymer is likely to localize at the blend interface. In addition, it has been shown that non-covalent and non-specific molecular interactions between polymers and MWNTs can lead to polymer adsorption which is partially irreversible, even during further mixing steps.²⁷ This strategy was further extended to drive the MWNTs to the interface of the blend by wrapping them with PS-r-PMMA in the first mixing step. To investigate the localization of MWNTs and PS-r-PMMA modified MWNTs in the blend, TEM (transmission electron microscopy) and STEM (scanning electron microscopy in transmission mode) were performed on blend samples which were annealed for 5 h at 220 °C and subsequently quenched to arrest the microstructure. The TEM and STEM images of various 85/15 PαMSAN/PMMA blends show a droplet–matrix structure after phase separation, both for the neat blend (Fig. 1a) as well as for the blends with MWNTs (Fig. 1b–d). The high resolution TEM image of the blend with MWNTs shows that the MWNTs are mostly localized in the grey PαMSAN matrix phase while exhibiting some agglomeration (Fig. 1b). In our earlier study, we have clearly demonstrated that MWNTs have a specific affinity towards the PαMSAN phase.¹⁶ Hence, in the absence of PS-r-PMMA the MWNTs were localized mostly in the PαMSAN phase as expected. In presence of PS-r-PMMA (Fig. 1c and d), the nanotubes are localized partially at the interface. This clearly shows that the localization of MWNTs can be tuned in binary blends by tailoring the mixing strategy and components to allow polymer wrapping of the MWNTs. Interestingly, in the case of blends with PS-r-PMMA modified MWNTs, an additional feature is worth noting. The MWNTs, owing to their large aspect ratio, are observed to bridge the droplets as indicated in the STEM images (Fig. 1d). This phenomenon can be quite interesting from electrical charge transport point of view and will be discussed later.

Fig. 1 TEM images showing phase separated (a) 85/15 PαMSAN/PMMA blends; (b) 85/15 PαMSAN/PMMA blends with 0.25 wt% MWNTs; STEM images showing (c) 85/15 PαMSAN/PMMA blends with PS-r-PMMA modified MWNTs; (d) high resolution image of (c). The concentration of MWNTs is 0.25 wt% in figures (b–d).

Although in both the blends containing MWNTs a droplet–matrix morphology was obtained after phase separation, there is a distinct difference. In the case of blends with PS-r-PMMA modified MWNTs, the PMMA droplets are significantly larger as compared to those in the blends with unmodified MWNTs. In addition, the former blends reveal a bimodal distribution of droplet sizes. It is imperative to mention that a coarser morphology in the presence of copolymer is counter intuitive and is rather opposite to the expected trend when using a copolymer as compatibilizer during melt mixing.^28–30 However, it should be kept in mind that in phase separating blends, the copolymer does not merely stabilize against coarsening but may also affect the thermodynamics and kinetics of phase separation.

Rheologically determined phase separation temperature and kinetics

To shed more light on the differences in morphology, we took a closer look at the evolution of the storage modulus during annealing by performing isochronal temperature sweep and time sweep experiments. The storage modulus provides indirect information about the microstructure since the presence of a droplet–matrix interface as well as the percolation of MWNTs in one phase due to their selective localization will provide additional elasticity. The strain amplitudes used in these experiments are chosen to be in the linear viscoelastic region. Hence, the dynamic measurements do not interfere with the thermodynamics of phase separation.

From the storage modulus G′ versus temperature curves, the rheological phase separation temperature can be determined (Fig. 2a). It is important to note that for blends with a large dynamic asymmetry and rather small component moduli, the interface contribution can be resolved easily. Hence, in this case rheological measurements are more sensitive than other characterization techniques such as optical methods to obtain the cloud point temperature. However, for blends with a small dynamic asymmetry and large component moduli, like the ones investigated here, the rheological method works only when the phase-separated structure is well developed, which occurs at high temperatures and after long annealing times. In light of this, the rheological transition temperature also depends on the kinetics of phase growth. The rheological phase separation temperatures in the case of blends with PS-r-PMMA modified MWNTs and MWNTs are observed to be similar to that of the neat blend (ca. 191 °C). The time sweep experiments (Fig. 2b) reveal that all blends more or less attain equilibrium after annealing for 5 h at 220 °C. However, the moduli are about an order of magnitude higher in the blends with PS-r-PMMA modified MWNTs as compared to the blends with only MWNTs. The latter indicates an improved percolation of the PS-r-PMMA modified MWNTs, which may result from their enhanced interfacial localization combined with droplet bridging.³¹ Unfortunately, this dominance of the elasticity of the MWNT network in the time evolution of the storage modulus also hampers gathering information about the coarsening process from rheology.

Fig. 2 (a) Storage modulus at 0.1 rad s⁻¹ versus temperature; (b) storage modulus at 0.1 rad s⁻¹ versus annealing time at 220 °C.

In summary, it is evident that the rheological phase separation temperature is nearly identical in all the blends investigated here and thus we can expect that all blends are at a similar quench depth when annealed at 220 °C for 5 h. Nevertheless, the STEM micrographs (Fig. 1c and d) demonstrate a bimodal droplet size distribution in the blends with PS-r-PMMA modified MWNTs whereas a uniform droplet size distribution with relatively small droplets occurs in the other blends (Fig. 1a and b). Clearly, the copolymer affects the coarsening process. It is worth mentioning here that the copolymer used in this study is an asymmetric random copolymer of PS and PMMA with a composition of 70 wt% styrene and 30 wt% methyl methacrylate. It has been reported that the structure of a copolymer is an important consideration in blending polymers.³² Previous studies have demonstrated that stronger interfaces can be achieved with long symmetric random copolymers as compared to long diblock copolymers or asymmetric random copolymers.³² It has also been suggested that when adding copolymers to an immiscible blend more effective lowering in interfacial tension may be achieved as compared to that reached with block copolymers.³³ Hence, block copolymers would be more effective in reducing the droplet size as compared to random copolymers. The effect of addition of co-polymer on the droplet size during the phase separation process in partially miscible blends has also been described.³⁴ Similar to immiscible blends, the blends with a symmetric copolymer have smaller domain sizes than the binary blend whereas the blend with the asymmetric co-polymer has a larger domain size than the binary blend. Recently, Marangoni stresses due to concentration gradients have also been shown to enhance coarsening.³¹ The random copolymer in our study is likely to increase such Marangoni stresses. It is also important to note that random copolymers (like the PS-r-PMMA used here) that preferentially localize at the interface usually encapsulate the droplet phase. When annealed for longer times (5 h in our case), the size of the encapsulated droplets has been observed to increase due to coalescence which is initiated by the overlap of the copolymer containing encapsulation regions around the droplets.²⁸ On the other hand, block copolymers suppress coalescence due to the steric hindrance caused by their presence at the interface.^35,36 Vinckier et al.³⁷ observed that the droplet size distribution in a phase separating PMMA/PαMSAN system became broader as a function of time and attributed this to Ostwald ripening. In this process, molecules of the dispersed PMMA phase diffuse from the small droplets to the larger ones, thereby resulting in an increase in droplet size. However, the numeric and volumetric mean droplet diameter did not increase to the same extent thereby increasing the polydispersity in droplet size. It can be anticipated that the presence of a random copolymer at the droplet interfaces affects this Ostwald ripening process thereby affecting the droplet size distribution.

In most polymer blends, the MWNTs were observed to localize in one of the phases and only a few studies report MWNTs at the interface.^18,27,29 It is important to note that unlike block co-polymers, the small functional groups on the surface of MWNTs cannot reptate the parent MWNT at the interface. Our STEM images suggest that the MWNTs were partially localized at the interface driven by the random co-polymer (PS-r-PMMA) which tends to encapsulate the droplets. Hence, the combination of above-mentioned factors is suggested to contribute to the larger and non-uniform droplet size in blends with PS-r-PMMA modified MWNTs. However, to pinpoint the exact mechanism, further research is needed, which is beyond the scope of our present interest.

Evolution of phase morphology: effect on electrical conductivity and EMI shielding

AC electrical conductivity at 220 °C was measured by using broadband dielectric spectroscopy. Before discussing the evolution of the AC electrical conductivity of the blends as a function of annealing time, it is important to evaluate the electrical conductivity in the components in presence of MWNTs. An electrical percolation threshold between 1–2 wt% MWNTs was observed in PMMA/MWNT composites whereas in the case of PαMSAN/MWNT composites we did not obtain any percolation even at 3 wt% MWNTs (results not shown). It is evident that the percolation threshold greatly depends on the matrix or more specifically on the polarity and surface free energy of the polymer. The surface free energy of the PαMSAN phase is known to be higher than that of PMMA¹⁶ and hence, matrix–filler interaction is expected to be more pronounced in PαMSAN/MWNT composites than in PMMA/MWNT composites. This explains the higher percolation threshold in PαMSAN/MWNT composites.

In order to study the evolution of the electrical conductivity in situ as a function of time, the blend samples were loaded between electrodes preheated at 220 °C and the AC electrical conductivity was measured in intervals of 30 min during a time span of 5 h. As expected 85/15 PαMSAN/PMMA blends with MWNTs showed no change in the electrical conductivity, even after annealing for 5 h (Fig. 3a). Recall that in these blends the MWNTs were localized mostly in the matrix phase (Fig. 1b). In addition, we did not observe any droplet-bridging effect in this case. Hence, the blends with MWNTs lack an interconnected network of MWNTs resulting in an insulating material. In our earlier study, 85/15 PαMSAN/PMMA blends with 2 wt% NH₂-MWNTs showed an AC electrical conductivity of 10⁻⁶ S cm⁻¹ after 5 h annealing at 220 °C. Interestingly, the 85/15 PαMSAN/PMMA blends with PS-r-PMMA modified MWNTs show a moderate electrical conductivity of 10⁻⁷ S cm⁻¹ immediately after loading the sample at 220 °C. It should be noted that the sample is at room temperature when being placed between the electrodes and heated till 220 °C inside the sample holder. Since it takes 5 min to reach 220 °C the very first stage of phase separation cannot be picked up in this way. The conductivity further increased by ca. 2 orders of magnitude upon annealing (Fig. 3b). It is envisaged that upon phase separation, the MWNTs migrate to the interface mediated by PS-r-PMMA. Further, as the blend coarsens the droplets grow bigger and the MWNTs bridge the droplets leading to an interconnected network like structure (Fig. 1d). This most likely resulted in high electrical conductivity at a rather low concentration of MWNTs. As discussed, a high electrical conductivity at a low concentration of conducting particles is in great demand. This helps in preserving the physical properties of the polymer which becomes brittle at higher filler loadings.

Fig. 3 Evolution of electrical conductivity (σ_AC) at 220 °C as a function of annealing time for 85/15 PαMSAN/PMMA blends with (a) MWNTs and (b) PS-r-PMMA modified MWNTs. The effective concentration of MWNTs is 0.25 wt% in both figures (a and b).

EMI shielding depends on the conductivity of the material since both reflection and absorption of EM waves increase with the conductivity of the material.^38,39 In order to further explore these composites for EMI shielding applications, we annealed the blends at 220 °C for 5 h to obtain phase separated structures (as in Fig. 1). The EMI shielding effectiveness of the 85/15 PαMSAN/PMMA blend with PS-r-PMMA modified MWNTs in the X-band and the K_u-band is depicted in Fig. 4. The neat 85/15 PαMSAN/PMMA blend (not shown here for clarity) and the blend with MWNTs are transparent to the EM radiation as can be observed from Fig. 4. Interestingly, in presence of PS-r-PMMA modified MWNTs, the 85/15 PαMSAN/PMMA blend shows a shielding effectiveness (SE) of −12 dB at 18 GHz. The enhanced SE is due to the percolating network of the MWNTs mediated by PS-r-PMMA. In order to evaluate the effect of the sample thickness on the EM shielding, specimens with different thicknesses were studied as shown in Fig. 4b and c. Fig. 4b illustrates the total shielding effectiveness for different shield thickness for 85/15 PαMSAN/PMMA blend with PS-r-PMMA modified MWNTs. It was observed that the total SE scaled with the thickness of the specimen. These results indicate that droplets covered with MWNTs can attenuate the incoming EM radiation effectively even at low concentration of the conducting filler.

Fig. 4 (a) EMI shielding effectiveness of 85/15 PαMSAN/PMMA blends after 5 h of annealing with a shield thickness of 6.6 mm, (b) shielding effectiveness of blend with PS-r-PMMA modified MWNTs as a function of frequency at different thicknesses, (c) shielding effectiveness of blend with PS-r-PMMA modified MWNTs as a function of thickness at 18 GHz.

Conclusions

85/15 PαMSAN/PMMA blends were melt mixed with PS-r-PMMA modified MWNTs and subsequently allowed to phase separate above their phase separation temperature (i.e. 220 °C). In presence of PS-r-PMMA, the nanotubes are localized partially at the interface. This clearly shows that the localization of MWNTs can be tuned in binary blends by tailoring the mixing strategy and components to allow polymer wrapping of the MWNTs. Our STEM images suggest that the MWNTs were partially localized at the interface driven by the random co-polymer (PS-r-PMMA) which tends to encapsulate the droplets resulting in larger and non-uniform droplet size in blends with PS-r-PMMA modified MWNTs. In addition, ‘droplet-bridging’ effect is observed which further resulted in high electrical conductivity at a relatively low fraction of MWNTs (0.25 wt%) in the blends. Moreover, a substantial EMI shielding effectiveness was recorded. We hypothesize that the incoming EM radiation suffers multiple scattering by the MWNTs distributed along the droplet interfaces and finally gets attenuated (see Scheme 1). It can be clearly observed that the SE obtained in this study is quite high especially considering the very low concentration of MWNTs. Further research is required to provide a more in depth understanding of this phenomenon. Such materials can further be explored for flexible EMI shielding materials.

Scheme 1 A cartoon illustrating the ‘droplet-bridging’ effect and the attenuation of incoming EM radiation through multiple scattering from the MWNT covered droplets in 85/15 PαMSAN/PMMA blends mediated by PS-r-PMMA.

Acknowledgements

R. Cardinaels is indebted to the Research Foundation Flanders (FWO) for a postdoctoral fellowship at KU Leuven.

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