Electroconductive Composites from Polystyrene Block Copolymers and Cu–Alumina Filler

Technological advancements and development of new materials may lead to the manufacture of sustainable energy-conducting devices used in the energy sector. This research attempts to fabricate novel electroconductive and mechanically stable nanocomposites via an electroless deposition (ELD) technique using electrically insulating materials. Metallic Cu is coated onto Al2O3 by ELD, and the prepared filler is then integrated (2–14 wt %) into a matrix of polystyrene-block-poly(ethylene-ran-butylene)-block-polystyrene-graft-maleic anhydride (PS-b-(PE-r-B)-b-PS-g-MA). Considerable variations in composite phases with filler inclusion exist. The Cu crystallite growth onto Al2O3 was evaluated by X-ray diffraction (XRD) analysis and energy dispersive spectrometry (EDS). Scanning electron microscopy (SEM) depicts a uniform Cu coating on Al2O3, while homogeneous filler dispersion is exhibited in the case of composites. The electrical behavior of composites is enhanced drastically (7.7 × 10−5 S/cm) upon incorporation of Cu–Al2O3 into an insulating polymer matrix (4.4 × 10−16 S/cm). Moreover, mechanical (Young’s modulus, tensile strength and % elongation at break) and thermal (thermogravimetric analysis (TGA), derivative thermogravimetry (DTG), and differential scanning calorimetry (DSC)) properties of the nanocomposites also improve substantially. These composites are likely to meet the demands of modern high-strength electroconductive devices.


Introduction
Technological advances highly depend on the development of a wide diversity of new materials. Conductive polymer composites (CPCs) have an array of applications in various industries, among them the electronic industry, which made revolutionary developments both in manufacturing and recycling. Electrostatic discharge (ESD) and electromagnetic interference (EMI) are phenomena that affect the economy of the electronic industry. They can arise during manufacturing, packing, conveyance, and working. Thus, the use of appropriate EMI-shielding materials to reduce electric energy losses is essential [1]. The ever-growing electronic waste (e-waste) is now posing devastating impact on the environment due to its accumulation. One way to reduce this accumulation is to increase the life span of electronics to protect them from the detrimental effects of EMI and ESD. Design and

Pretreatment of Al 2 O 3
Pretreatment of Al 2 O 3 was done before the deposition step, followed by surface cleaning and surface activation. To avoid tedious filtration steps, Al 2 O 3 was packed in commercially available silk cloth (160 mesh) and was dipped in subsequent solutions rather than dispersion in solution, which may also increase the reaction time. First, the Al 2 O 3 substrate was dipped in concentrated HNO 3 (2 min) to remove oil and dirt. Acid-cleaned Al 2 O 3 was dipped in catalytic activator solution containing 0.03 mmol of PdCl 2 and 0.246 mmol of SnCl 2 in 40 mL of concentrated HCl (14 min). After activation, the substrate was introduced to a reduction bath made of 4.74 mmol of (C 4 H 10 BN) and 4.52 mmol of (H 3 BO 3 ) in a sufficient quantity of distilled water (7 min). Each step of pretreatment was followed by 1 min rinsing in distilled water. Pretreated Al 2 O 3 was then used for ELD of Cu.

Cu Coating on Pretreated Al 2 O 3
Pretreated Al 2 O 3 was dipped in an electroless plating bath (Table 1). After deposition, Cu-deposited Al 2 O 3 was rinsed with distilled water and oven-dried for 3-4 h at 40 • C. The prepared Cu-Al 2 O 3 powder was used further as conductive filler for insulating polymer matrix.

Synthesis of Conductive Composites
Conductive polymer composites (CPCs) were prepared by incorporating Cu-Al 2 O 3 filler with varied content (2,4,6,8,10,12, and 14 wt %) in PS-b-(PE-r-B)-b-PS-g-MA polymer matrix. PS-b-(PE-r-B)-b-PS-g-MA was dissolved in 30 mL chloroform followed by addition of the filler. The polymer-filler solution was stirred for 1-2 h at 750 rpm, poured into a Petri dish for film casting, and then detached from the mold after solvent evaporation. The prepared composite films were utilized for characterization. To attain accuracy in performance and results, samples were prepared in triplicates and the mean values were reported after characterization.

Morphological Analysis
The surface morphologies of pristine Al 2 O 3 , Cu-coated Al 2 O 3 , host polymer, and its respective composites were analyzed with SEM, obtained by a HT-Phys-UAJK microscope equipped with a secondary electron (SE) detector at 25 kV accelerating voltage. Fractured surfaces of composites were also examined by MIRA3 TESCAN (Nova 400 Nano, Salem, OR, USA) (SE detector at an accelerating voltage of 10 kV) to analyze the dispersion of filler in polymer matrix.

Energy Dispersive Spectrometry (EDS) Analysis
Elemental composition and atomic weight % of Cu coated Al 2 O 3 , host polymer and its respective composites were investigated by using a JSM6490LV (JEOL) microscope (Tokyo, Japan). The instrument was equipped with QUANTAX EDS XFlash detector 4010-Bruker (Billerica, MA, USA) at an accelerating voltage of 20 kV.

Analysis of Surface/Volume Resistivity and Electrical Conductivity
The surface resistivity (Ω/ ) and volume resistivity (Ω·cm) of composites were measured by the 4-probe method using a high-resistance meter by applying the ASTM D-257 test method [15] at room temperature. A 500 V direct current field was applied through electrodes made up of tungsten carbide. Since electrical conductivity is inversely proportional to volume resistivity, electrical conductivities (S/cm) were calculated as the inverse of the volume resistivities (Ω·cm).

Analysis of Mechanical Properties
The mechanical features of composites were examined by calculating the Young's modulus (MPa), tensile strength (MPa), and % elongation at break according to the ASTM D638-02 [16] and ASTM D638-03 [17] test procedures for mechanical analysis. An Instron tester (4465UK, Norwood, MA, USA) was used at 20 ± 2 • C by subjecting samples with dimensions of 0.8-1.0 mm thickness and 6 mm× 70 mm (width × gauze length).

Analysis of Thermal Properties
Thermogravimetric analysis (TGA) was carried out with a Perkin Elmer TGA-7 (Waltham, MA, USA) in the 50-550 • C temperature range at 20 • C/min in dynamic atmosphere (20 mL/min N 2 flow) using a 2 mg sample. Non-isothermal conditions were used for recording thermal analytical results. A DSC 404-NETZSCH instrument was used for differential scanning calorimetry (DSC) analysis in the 20-500 • C range at 20 • C/min.

XRD Analysis of Pristine and Cu-Coated Al 2 O 3 Powder and Composite Films
The prepared filler was analyzed with XRD for the determination of phase change and particle size. Scherrer's equation is used to calculate the particle size of pristine and Cu-coated Al 2 O 3 as expressed in Equation (1) [18]: where D = crystallite size (nm); λ = wavelength; K = Scherrer's constant; β = angular width (radians); and θ = Bragg's angle.
Interplanar spacing between atoms within the crystallite structure is denoted by d-spacing. Bragg's equation used for the determination of d-spacing of pristine and Cu-coated Al 2 O 3 is given in Equation (2).
The XRD spectrum of pristine Al 2 O 3 powder is presented in Figure 1. The XRD spectrum of pristine Al2O3 powder is presented in Figure 1. The three peaks at 13.13°, 46.04°, and 67.28° (2θ values), correspond to the Al2O3 phase. The strongest diffraction peak is observed at 67.28° with minimum d-spacing 0.139 nm. At 46.04°, another peak exists corresponding to Al2O3 with d-spacing 0.197 nm. The average size of the cubic lattice of Al2O3 is approximately 8.8 nm. The XRD spectrum of Cu-Al2O3 ( Figure 2) depicts the strongest peak at 43.5° and a relatively less intense peak at 50.6° corresponding to the (111) and (200) lattice planes of Cu, respectively. The strongest diffraction peak at 43.5° is characteristic of a face-centered cubic structure with d-spacing of 0.21 nm; this confirms deposition of crystal-structured metallic Cu on the substrate [18,19]. The relative peak intensity at 2θ = 67.3° clearly represents the XRD pattern of pristine Al2O3 (Figure 1), whose amount was lower in the composite material. The disturbance observed in the peak corresponding to the Al2O3 phase is due to the change in the nature of original Al2O3 after the deposition of Cu. The average crystallite size of Cu-Al2O3 was calculated as approximately 26.2 nm. XRD analysis also revealed that average crystallite size of Al2O3 increased from 8.8 nm to 26.2 nm, which confirms the deposition of Cu crystallites, with an increase in the mean thickness to ~17.4 nm. A similar XRD pattern was reported in literature [20,21], where the strongest peak of electroless deposited-Cu appeared at 2θ = 43°. For determining the effect of Cu-Al2O3 filler in the host polymer matrix, XRD spectra of the polymer with 2 wt % and 14 wt % of Cu-Al2O3-loading were recorded ( Figure 3). The XRD pattern of neat PS-b-(PE-r-B)-b-PS-g-MA shows a broad peak at 10°-27° and one relatively less intense peak at 48.9°, which confirms its amorphous structure. Upon 2 wt % Cu-Al2O3 loading in PS-b-(PE-r-B)-b-PS-g-MA, two peaks at 42.6°and 49.9° are observed. The peak at 42.6° is attributed to the crystalline The XRD spectrum of Cu-Al 2 O 3 ( Figure 2) depicts the strongest peak at 43.5 • and a relatively less intense peak at 50.6 • corresponding to the (111) and (200) lattice planes of Cu, respectively. The strongest diffraction peak at 43.5 • is characteristic of a face-centered cubic structure with d-spacing of 0.21 nm; this confirms deposition of crystal-structured metallic Cu on the substrate [18,19]. The relative peak intensity at 2θ = 67.3 • clearly represents the XRD pattern of pristine Al 2 O 3 ( Figure 1), whose amount was lower in the composite material. The disturbance observed in the peak corresponding to the Al 2 O 3 phase is due to the change in the nature of original Al 2 O 3 after the deposition of Cu. The average crystallite size of Cu-Al 2 O 3 was calculated as approximately 26.2 nm. XRD analysis also revealed that average crystallite size of Al 2 O 3 increased from 8.8 nm to 26.2 nm, which confirms the deposition of Cu crystallites, with an increase in the mean thickness to~17.4 nm. A similar XRD pattern was reported in literature [20,21], where the strongest peak of electroless deposited-Cu appeared at 2θ = 43 • . The XRD spectrum of pristine Al2O3 powder is presented in Figure 1. The three peaks at 13.13°, 46.04°, and 67.28° (2θ values), correspond to the Al2O3 phase. The strongest diffraction peak is observed at 67.28° with minimum d-spacing 0.139 nm. At 46.04°, another peak exists corresponding to Al2O3 with d-spacing 0.197 nm. The average size of the cubic lattice of Al2O3 is approximately 8.8 nm. The XRD spectrum of Cu-Al2O3 ( Figure 2) depicts the strongest peak at 43.5° and a relatively less intense peak at 50.6° corresponding to the (111) and (200) lattice planes of Cu, respectively. The strongest diffraction peak at 43.5° is characteristic of a face-centered cubic structure with d-spacing of 0.21 nm; this confirms deposition of crystal-structured metallic Cu on the substrate [18,19]. The relative peak intensity at 2θ = 67.3° clearly represents the XRD pattern of pristine Al2O3 (Figure 1), whose amount was lower in the composite material. The disturbance observed in the peak corresponding to the Al2O3 phase is due to the change in the nature of original Al2O3 after the deposition of Cu. The average crystallite size of Cu-Al2O3 was calculated as approximately 26.2 nm. XRD analysis also revealed that average crystallite size of Al2O3 increased from 8.8 nm to 26.2 nm, which confirms the deposition of Cu crystallites, with an increase in the mean thickness to ~17.4 nm. A similar XRD pattern was reported in literature [20,21], where the strongest peak of electroless deposited-Cu appeared at 2θ = 43°. For determining the effect of Cu-Al2O3 filler in the host polymer matrix, XRD spectra of the polymer with 2 wt % and 14 wt % of Cu-Al2O3-loading were recorded ( Figure 3). The XRD pattern of neat PS-b-(PE-r-B)-b-PS-g-MA shows a broad peak at 10°-27° and one relatively less intense peak at 48.9°, which confirms its amorphous structure. Upon 2 wt % Cu-Al2O3 loading in PS-b-(PE-r-B)-b-PS-g-MA, two peaks at 42.6°and 49.9° are observed. The peak at 42.6° is attributed to the crystalline For determining the effect of Cu-Al 2 O 3 filler in the host polymer matrix, XRD spectra of the polymer with 2 wt % and 14 wt % of Cu-Al 2 O 3 -loading were recorded ( Figure 3). The XRD pattern of neat PS-b-(PE-r-B)-b-PS-g-MA shows a broad peak at 10 • -27 • and one relatively less intense peak at 48.9 • , which confirms its amorphous structure. Upon 2 wt % Cu-Al 2 O 3 loading in PS-b-(PE-r-B)-b-PS-g-MA, two peaks at 42.6 • and 49.9 • are observed. The peak at 42.6 • is attributed to the crystalline nature of Cu, while another peak at 49.9 • might represent a slight peak shift from 48.92 • corresponding to amorphous phase of the polymer. At 14 wt % Cu-Al 2 O 3 loading, three peaks at 2θ = 36.2 • , 42.9 • and 50.1 • arise. The sharp peak observed at 42.9 • is characteristic of metallic Cu inclusion supported by another peak at 50.1 • and thus confirms the crystalline phase of the prepared composites.
nature of Cu, while another peak at 49.9° might represent a slight peak shift from 48.92° corresponding to amorphous phase of the polymer. At 14 wt % Cu-Al2O3 loading, three peaks at 2θ = 36.2°, 42.9° and 50.1° arise. The sharp peak observed at 42.9° is characteristic of metallic Cu inclusion supported by another peak at 50.1° and thus confirms the crystalline phase of the prepared composites.

Morphological Study of Cu-Al2O3 Filler and Block Copolymer Composites
SEM analysis was used to determine the surface morphology and crystalline structure of the materials. An SEM micrograph of pristine Al2O3 and Cu-Al2O3 powder are shown in Figure 4a,b, respectively, showing a uniform Cu coating on the alumina surface. The dispersion of the filler is improved as compared to pristine alumina powder. The Cu-coated Al2O3 particles exhibit fine-scale roughness, characteristic of metal coating [22,23]. Silvain and co-workers also deposited Cu onto submicron-sized Al2O3 particles [24]. Their work revealed uniform and fine coating of metallic Cu and increased average particle size of Al2O3 particles after Cu deposition. The SEM image of Cu-Al2O3 from Wang and co-workers [25] showed good similarity ( Figure 4c). In comparison, Krupa and co-workers deposited Ag on polyimide particles ( Figure 4d) [26]. In all these cases, the ELD-plating technique was used. The morphological properties looked similar and seem to be rather irrespective of the type of substrate.

Morphological Study of Cu-Al 2 O 3 Filler and Block Copolymer Composites
SEM analysis was used to determine the surface morphology and crystalline structure of the materials. An SEM micrograph of pristine Al 2 O 3 and Cu-Al 2 O 3 powder are shown in Figure 4a,b, respectively, showing a uniform Cu coating on the alumina surface. The dispersion of the filler is improved as compared to pristine alumina powder. The Cu-coated Al 2 O 3 particles exhibit fine-scale roughness, characteristic of metal coating [22,23]. Silvain and co-workers also deposited Cu onto submicron-sized Al 2 O 3 particles [24]. Their work revealed uniform and fine coating of metallic Cu and increased average particle size of Al 2 O 3 particles after Cu deposition. The SEM image of Cu-Al 2 O 3 from Wang and co-workers [25] showed good similarity ( Figure 4c). In comparison, Krupa and co-workers deposited Ag on polyimide particles ( Figure 4d) [26]. In all these cases, the ELD-plating technique was used. The morphological properties looked similar and seem to be rather irrespective of the type of substrate.

Morphological Study of Cu-Al2O3 Filler and Block Copolymer Composites
SEM analysis was used to determine the surface morphology and crystalline structure of the materials. An SEM micrograph of pristine Al2O3 and Cu-Al2O3 powder are shown in Figure 4a,b, respectively, showing a uniform Cu coating on the alumina surface. The dispersion of the filler is improved as compared to pristine alumina powder. The Cu-coated Al2O3 particles exhibit fine-scale roughness, characteristic of metal coating [22,23]. Silvain and co-workers also deposited Cu onto submicron-sized Al2O3 particles [24]. Their work revealed uniform and fine coating of metallic Cu and increased average particle size of Al2O3 particles after Cu deposition. The SEM image of Cu-Al2O3 from Wang and co-workers [25] showed good similarity ( Figure 4c). In comparison, Krupa and co-workers deposited Ag on polyimide particles ( Figure 4d) [26]. In all these cases, the ELD-plating technique was used. The morphological properties looked similar and seem to be rather irrespective of the type of substrate.  With the incorporation of lowest filler content (2 wt %), homogenous dispersion of filler is observed in both the composites predicting good filler-polymer interaction. At least 10 wt % filler is required to observe the initiation of particleto-particle connectivities, which improve throughout the matrix when the filler content is further increased to 14 wt %. The comparison of Figure 5b,e illustrates the decreased interparticle distance. The shiny small areas in the SEM images resemble the presence of the metal coated on ceramic filler. Even smaller interparticle distance could be achieved with filler loadings higher than 14 wt %, but this compromises the mechanical performance considering the properties of ceramics. A clear transition in the particle shape and surface roughness takes place upon Cu metallization. The uniform growth of Cu crystallites on Al2O3 explains the change in morphology regarding particle distribution, which ultimately affects the mean coating thickness. At higher filler loading, agglomerates or islands of the filler particles are formed within the matrix material, which helps the smooth transfer of electrons. Individual nanosized filler particles are not distinctly visible in SEM micrographs because of this phenomenon [27]. With the incorporation of lowest filler content (2 wt %), homogenous dispersion of filler is observed in both the composites predicting good filler-polymer interaction. At least 10 wt % filler is required to observe the initiation of particle-to-particle connectivities, which improve throughout the matrix when the filler content is further increased to 14 wt %. The comparison of Figure 5b,e illustrates the decreased interparticle distance. The shiny small areas in the SEM images resemble the presence of the metal coated on ceramic filler. Even smaller interparticle distance could be achieved with filler loadings higher than 14 wt %, but this compromises the mechanical performance considering the properties of ceramics. A clear transition in the particle shape and surface roughness takes place upon Cu metallization. The uniform growth of Cu crystallites on Al 2 O 3 explains the change in morphology regarding particle distribution, which ultimately affects the mean coating thickness. At higher filler loading, agglomerates or islands of the filler particles are formed within the matrix material, which helps the smooth transfer of electrons. Individual nanosized filler particles are not distinctly visible in SEM micrographs because of this phenomenon [27]. With the incorporation of lowest filler content (2 wt %), homogenous dispersion of filler is observed in both the composites predicting good filler-polymer interaction. At least 10 wt % filler is required to observe the initiation of particleto-particle connectivities, which improve throughout the matrix when the filler content is further increased to 14 wt %. The comparison of Figure 5b,e illustrates the decreased interparticle distance. The shiny small areas in the SEM images resemble the presence of the metal coated on ceramic filler. Even smaller interparticle distance could be achieved with filler loadings higher than 14 wt %, but this compromises the mechanical performance considering the properties of ceramics. A clear transition in the particle shape and surface roughness takes place upon Cu metallization. The uniform growth of Cu crystallites on Al2O3 explains the change in morphology regarding particle distribution, which ultimately affects the mean coating thickness. At higher filler loading, agglomerates or islands of the filler particles are formed within the matrix material, which helps the smooth transfer of electrons. Individual nanosized filler particles are not distinctly visible in SEM micrographs because of this phenomenon [27].

EDS Analysis of Cu-Al2O3 Filler and Block Copolymer Composites
EDS was used to study the elemental composition of Cu-Al2O3 filler and Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA composites (Table 2). Cu, Al, and Pd were detected in the filler material. The high content of Cu (67.7%) followed by Al (30.4%) and Pd (1.9%) confirms the effective deposition of Cu-Al2O3 via the ELD process. Pd was present in small quantities, as it was used at a minor concentration for surface activation of the Al2O3 substrate.

Surface/Volume Resistivity and Electrical Conductivity of Block Copolymer Composites
The volume resistivity is the reciprocal of the electric conductivity. Measurement of the resistance across the materials' surface, which is in contact with the electrodes, is termed surface resistivity (Ω/sq or Ω/□) [28,29], while electrical resistance through a cube of insulating material is considered as volume resistivity (Ω·cm). The host matrix polymers are usually non-conducting in nature and contain an insignificant number of charge carriers in free-state. Thus, the electrical properties of such matrix polymer composites almost exclusively depend on the selection of filler and its ability to form smooth conductive networks throughout the matrix [30][31][32][33][34][35]. The surface resistivity of Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA matrix composites with increasing filler loadings (2-14 wt %) were studied. The surface and volume resistivity of neat polymer was also analyzed to determine its electrical behavior as intrinsic or extrinsic conducting polymer matrix. Table 3 shows the surface resistivity, the volume resistivity, and the electrical conductivity. The values for PS-b-(PE-r-B)-b-PSg-MA are 2.30 × 10 14 Ω/□, 2.3 × 10 15 Ω·cm, and 4.348 × 10 −16 S/cm, respectively, which confirms that it cannot act as intrinsic conducting polymer; although bulky aromatic rings are present as pendants, the main chain is saturated, rendering an insulation material. Table 3. Surface/volume resistivity and electrical conductivity of Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA composites.   (Table 2). Cu, Al, and Pd were detected in the filler material. The high content of Cu (67.7%) followed by Al (30.4%) and Pd (1.9%) confirms the effective deposition of Cu-Al 2 O 3 via the ELD process. Pd was present in small quantities, as it was used at a minor concentration for surface activation of the Al 2 O 3 substrate.

Surface/Volume Resistivity and Electrical Conductivity of Block Copolymer Composites
The volume resistivity is the reciprocal of the electric conductivity. Measurement of the resistance across the materials' surface, which is in contact with the electrodes, is termed surface resistivity (Ω/sq or Ω/ ) [28,29], while electrical resistance through a cube of insulating material is considered as volume resistivity (Ω·cm). The host matrix polymers are usually non-conducting in nature and contain an insignificant number of charge carriers in free-state. Thus, the electrical properties of such matrix polymer composites almost exclusively depend on the selection of filler and its ability to form smooth conductive networks throughout the matrix [30][31][32][33][34][35]. The surface resistivity of Cu-Al 2 O 3 /PS-b-(PE-r-B)-b-PS-g-MA matrix composites with increasing filler loadings (2-14 wt %) were studied. The surface and volume resistivity of neat polymer was also analyzed to determine its electrical behavior as intrinsic or extrinsic conducting polymer matrix. Table 3 shows the surface resistivity, the volume resistivity, and the electrical conductivity. The values for PS-b-(PE-r-B)-b-PS-g-MA are 2.30 × 10 14 Ω/ , 2.3 × 10 15 Ω·cm, and 4.348 × 10 −16 S/cm, respectively, which confirms that it cannot act as intrinsic conducting polymer; although bulky aromatic rings are present as pendants, the main chain is saturated, rendering an insulation material.
With the inclusion of small amounts of filler (2 wt %), the surface resistivity of the corresponding composite readily drops from insulating to antistatic region. The corresponding electrical conductivity increases to 2.381 × 10 14 Ω·cm. This immediate shift from insulating to antistatic region might be attributed to the connection with unsaturated side chain substitutions like maleic anhydride and benzene groups, which essentially help to enhance particle-to-particle interaction [36]. By increasing the loading of Cu-Al 2 O 3 filler in the polymer matrix from 4 to 12 wt %, the surface and volume resistivity drop from 5.8 × 10 9 to 1.3 × 10 8 Ω/ and from 4.2 × 10 13 to 4.5 × 10 4 Ω·cm, respectively. This drop shifts the conductive properties of the material from the antistatic to the static dissipative region [37]. Upon further incorporation of filler (14 wt %), the surface resistivity drops drastically to 4.0 × 10 4 Ω/ , while the volume resistivity and electric conductivity changed only substantially compared to 12 wt % filler loading. The gradual increment in conductivity with addition of 2-14 wt % filler is shown in Figure 6. It confirms network formation as suggested by the SEM results, showing the transition from an insulating to a semiconducting region. Table 3. Surface/volume resistivity and electrical conductivity of Cu-Al 2 O 3 /PS-b-(PE-r-B)-b-PSg-MA composites.

Cu-Al 2 O 3 (wt %)
Surface With the inclusion of small amounts of filler (2 wt %), the surface resistivity of the corresponding composite readily drops from insulating to antistatic region. The corresponding electrical conductivity increases to 2.381 × 10 14 Ω·cm. This immediate shift from insulating to antistatic region might be attributed to the connection with unsaturated side chain substitutions like maleic anhydride and benzene groups, which essentially help to enhance particle-to-particle interaction [36]. By increasing the loading of Cu-Al2O3 filler in the polymer matrix from 4 to 12 wt %, the surface and volume resistivity drop from 5.8 × 10 9 to 1.3 × 10 8 Ω/□ and from 4.2 × 10 13 to 4.5 × 10 4 Ω·cm, respectively. This drop shifts the conductive properties of the material from the antistatic to the static dissipative region [37]. Upon further incorporation of filler (14 wt %), the surface resistivity drops drastically to 4.0 × 10 4 Ω/□, while the volume resistivity and electric conductivity changed only substantially compared to 12 wt % filler loading. The gradual increment in conductivity with addition of 2-14 wt % filler is shown in Figure 6. It confirms network formation as suggested by the SEM results, showing the transition from an insulating to a semiconducting region. The electron transfer responsible for conductivity throughout the Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA matrix takes place when interaction zones between filler and matrix find connections (Figure 7), establishing a web [38][39][40][41]. Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA composites are cost-effective materials, as they showed enhanced electrical conductivity and they are easy to prepare compared to previously cited literature [42]. Beyond the critical concentration or percolation limit, there is no further significant increase in electrical conductivity even though more filler is contained in the composite material. Once the saturation point is attained, further increase in filler loading may only increase the sum of conductive networks and does not contribute in further conductivity increments. In contrast, shielding effectiveness may increase when higher filler loadings are used [43][44][45].  The electron transfer responsible for conductivity throughout the Cu-Al 2 O 3 /PS-b-(PE-r-B)-b-PS-g-MA matrix takes place when interaction zones between filler and matrix find connections (Figure 7), establishing a web [38][39][40][41]. Cu-Al 2 O 3 /PS-b-(PE-r-B)-b-PS-g-MA composites are cost-effective materials, as they showed enhanced electrical conductivity and they are easy to prepare compared to previously cited literature [42]. Beyond the critical concentration or percolation limit, there is no further significant increase in electrical conductivity even though more filler is contained in the composite material. Once the saturation point is attained, further increase in filler loading may only increase the sum of conductive networks and does not contribute in further conductivity increments. In contrast, shielding effectiveness may increase when higher filler loadings are used [43][44][45]. With the inclusion of small amounts of filler (2 wt %), the surface resistivity of the corresponding composite readily drops from insulating to antistatic region. The corresponding electrical conductivity increases to 2.381 × 10 14 Ω·cm. This immediate shift from insulating to antistatic region might be attributed to the connection with unsaturated side chain substitutions like maleic anhydride and benzene groups, which essentially help to enhance particle-to-particle interaction [36]. By increasing the loading of Cu-Al2O3 filler in the polymer matrix from 4 to 12 wt %, the surface and volume resistivity drop from 5.8 × 10 9 to 1.3 × 10 8 Ω/□ and from 4.2 × 10 13 to 4.5 × 10 4 Ω·cm, respectively. This drop shifts the conductive properties of the material from the antistatic to the static dissipative region [37]. Upon further incorporation of filler (14 wt %), the surface resistivity drops drastically to 4.0 × 10 4 Ω/□, while the volume resistivity and electric conductivity changed only substantially compared to 12 wt % filler loading. The gradual increment in conductivity with addition of 2-14 wt % filler is shown in Figure 6. It confirms network formation as suggested by the SEM results, showing the transition from an insulating to a semiconducting region. The electron transfer responsible for conductivity throughout the Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA matrix takes place when interaction zones between filler and matrix find connections (Figure 7), establishing a web [38][39][40][41]. Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA composites are cost-effective materials, as they showed enhanced electrical conductivity and they are easy to prepare compared to previously cited literature [42]. Beyond the critical concentration or percolation limit, there is no further significant increase in electrical conductivity even though more filler is contained in the composite material. Once the saturation point is attained, further increase in filler loading may only increase the sum of conductive networks and does not contribute in further conductivity increments. In contrast, shielding effectiveness may increase when higher filler loadings are used [43][44][45].

Mechanical Properties of Cu-Al 2 O 3 /PS-b-(PE-r-B)-b-PS-g-MA Composites
Mechanical performance of a polymer matrix composite can be influenced by the composition and interaction of filler and matrix materials used. Geometrical aspects, such as structure shape and size of reinforcement material, considerably affect the mechanical behavior of composites [46]. For the synthesis of structurally resilient composites, filler dispersion and declustering is a prerequisite. Thus, by critically controlling the volume fraction of filler, mechanical properties were measured to analyze the effect of filler inclusion and to prevent any deterioration in mechanical properties of composites [45]. The mechanical behavior of Cu-Al 2 O 3 /polymer composites was examined by calculating Young's modulus, tensile strength, and % elongation at break of the composites with increasing filler loading (0-14 wt %).

Young's Modulus
Young's modulus is a quantitative parameter for the stiffness determination of elastic materials. It is defined as the ratio of applied stress to the strain along the same axis. The applied stress should be in the range in which Hook's law holds properly [47]. Young's modulus of neat block copolymer is 50 ± 3 MPa, which increased to 150 ± 3 MPa (Figure 8

Mechanical Properties of Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA Composites
Mechanical performance of a polymer matrix composite can be influenced by the composition and interaction of filler and matrix materials used. Geometrical aspects, such as structure shape and size of reinforcement material, considerably affect the mechanical behavior of composites [46]. For the synthesis of structurally resilient composites, filler dispersion and declustering is a prerequisite. Thus, by critically controlling the volume fraction of filler, mechanical properties were measured to analyze the effect of filler inclusion and to prevent any deterioration in mechanical properties of composites [45]. The mechanical behavior of Cu-Al2O3/polymer composites was examined by calculating Young's modulus, tensile strength, and % elongation at break of the composites with increasing filler loading (0-14 wt %).

Young's Modulus
Young's modulus is a quantitative parameter for the stiffness determination of elastic materials. It is defined as the ratio of applied stress to the strain along the same axis. The applied stress should be in the range in which Hook's law holds properly [47]. Young's modulus of neat block copolymer is 50 ± 3 MPa, which increased to 150 ± 3 MPa (Figure 8) with the gradual addition of reinforcement material. This gradual and constant increase in Young's modulus of composites with increased filler loading indicates enhancement in stiffness imparted by Al2O3.

Tensile Strength
The maximum stress that a material can endure before failing or breaking is known as tensile strength [48]. The incorporation of Cu-Al2O3 in the polymer matrix increases the tensile strength of the resultant composites. At 14 wt % Cu-Al2O3 loading, the tensile strength of the composite reached 82 ± 3 MPa, as compared to 15 ± 3 MPa of the neat polymer. Figure 9 shows the gradual increase of tensile strength with filler loading. Tensile strength is strongly dependent upon interfacial adhesion/bonding between filler and matrix and is aided by uniform filler dispersion. Interfacial adhesion determines the strength of such composites. The results suggest good compatibility between particulate filler and polymer matrix and confirms active transfer of stress from matrix to particulate filler [49][50][51]. The PS-b-(PE-r-B)-b-PS-g-MA/Cu-Al2O3 composites offer good strength and mechanical resistance, compared to previously reported polymer/metal-coated polymers [26], polymer/carbon [42], polymer/ceramic [52], polymer/mineral ( [53,54], ethylene-propylene-diene monomer rubber/Mg(OH)2) [55] and polymer/polymer composites [53], as illustrated in Table 4.

Tensile Strength
The maximum stress that a material can endure before failing or breaking is known as tensile strength [48]. The incorporation of Cu-Al 2 O 3 in the polymer matrix increases the tensile strength of the resultant composites. At 14 wt % Cu-Al 2 O 3 loading, the tensile strength of the composite reached 82 ± 3 MPa, as compared to 15 ± 3 MPa of the neat polymer. Figure 9 shows the gradual increase of tensile strength with filler loading. Tensile strength is strongly dependent upon interfacial adhesion/bonding between filler and matrix and is aided by uniform filler dispersion. Interfacial adhesion determines the strength of such composites. The results suggest good compatibility between particulate filler and polymer matrix and confirms active transfer of stress from matrix to particulate filler [49][50][51]. The PS-b-(PE-r-B)-b-PS-g-MA/Cu-Al 2 O 3 composites offer good strength and mechanical resistance, compared to previously reported polymer/metal-coated polymers [26], polymer/carbon [42], polymer/ceramic [52], polymer/mineral ( [53,54], ethylene-propylene-diene monomer rubber/Mg(OH) 2 ) [55] and polymer/polymer composites [53], as illustrated in Table 4.

Elongation at Break
Elongation at break is a quantitative parameter for the ductility of the material. It is defined as the percentage of elongation of a material from zero stress to the breaking point of that material [56]. The elongation at break is also an indicator for determining the toughness of two phase materials [57]. The elongation at break calculated for PS-b-(PE-r-B)-b-PS-g-MA polymer was 16.9% ± 0.4%. Cu-Al2O3/PS-b-(PE-r-B)-b-PS-g-MA composites with increasing filler loadings (0%-14%) showed a gradual decrease from 16.9% to 10.1% ( Figure 10). Polymers are ductile in nature while ceramics exhibit brittle behavior. Thus, the gradual increase in brittle behavior is due to the incorporation of the reinforcement material [58], and may arise from interstructural progression in which filler particles are dispersed in the interaggregate space [48]. At low filler loading, the matrix is not adequately reinforced. So, it could not withstand high load, and eventually failure happens at lower elongation. However, at higher filler loading, the matrix is increasingly reinforced and endures high load before the breaking point is reached. The reinforcement mechanism preludes that, at higher filler loading, the molecular mobility drops because of the formation of physical bonds among particles of filler and polymer molecule chains [43].

Elongation at Break
Elongation at break is a quantitative parameter for the ductility of the material. It is defined as the percentage of elongation of a material from zero stress to the breaking point of that material [56]. The elongation at break is also an indicator for determining the toughness of two phase materials [57]. The elongation at break calculated for PS-b-(PE-r-B)-b-PS-g-MA polymer was 16.9% ± 0.4%. Cu-Al 2 O 3 /PS-b-(PE-r-B)-b-PS-g-MA composites with increasing filler loadings (0%-14%) showed a gradual decrease from 16.9% to 10.1% ( Figure 10). Polymers are ductile in nature while ceramics exhibit brittle behavior. Thus, the gradual increase in brittle behavior is due to the incorporation of the reinforcement material [58], and may arise from interstructural progression in which filler particles are dispersed in the interaggregate space [48]. At low filler loading, the matrix is not adequately reinforced. So, it could not withstand high load, and eventually failure happens at lower elongation. However, at higher filler loading, the matrix is increasingly reinforced and endures high load before the breaking point is reached. The reinforcement mechanism preludes that, at higher filler loading, the molecular mobility drops because of the formation of physical bonds among particles of filler and polymer molecule chains [43].

Thermal Characteristics of Block Copolymer Composites
3.6.1. Thermogravimetric Analysis (TGA) TGA examines the thermal properties as the weight alteration upon heating during the phases of thermal breakdown. The thermal behavior determines the possible specific application fields of nanocomposites [59]. TGA thermograms of neat PS-b-(PE-r-B)-b-PS-g-MA and Cu-Al2O3 loaded composites from 0, 2, and 14 wt % are shown in Figure 11. A two-phase decomposition is observed for neat block copolymer. A slight dip at 250 °C indicates the presence of some residual low molecular weight compounds in the polymer. In the present conditions, the polymer remains stable up to 397 °C (8% weight loss). The second phase of decomposition starts at 397 °C (Tmax) and continues up to a final degradation temperature of 480 °C (99% weight loss at Tf). With the inclusion of 2 wt % Cu-Al2O3, the thermal stability of the composite is improved, where Tmax raises from 397 to 405 °C and Tf from 480 to 492 °C. At 14 wt % filler loading, Tmax and Tf are respectively 30 °C and 9 °C higher compared to the neat polymer. At this point (Tf 489 °C), 67% residue is still left. Upon heating the polymer, the long chains break down into small fragments which might have interacted with Cu-Al2O3 particles and got trapped into filler particles difficult to be decomposed further, thus improving the thermal stability of the PS-b-(PE-r-B)-b-PS-g-MA composites [56]. Similar behavior was observed previously, where thermal stability was enhanced due to filler incorporation which hindered the segmental movement of polymer when intermingled with small chains of the host polymer [27,48]. Analogous degradation patterns are seen in the derivative thermogravimetry (DTG) curves of host polymer and its composites ( Figure 12).  TGA examines the thermal properties as the weight alteration upon heating during the phases of thermal breakdown. The thermal behavior determines the possible specific application fields of nanocomposites [59]. TGA thermograms of neat PS-b-(PE-r-B)-b-PS-g-MA and Cu-Al 2 O 3 loaded composites from 0, 2, and 14 wt % are shown in Figure 11. A two-phase decomposition is observed for neat block copolymer. A slight dip at 250 • C indicates the presence of some residual low molecular weight compounds in the polymer. In the present conditions, the polymer remains stable up to 397 • C (8% weight loss). The second phase of decomposition starts at 397 • C (T max ) and continues up to a final degradation temperature of 480 • C (99% weight loss at T f ). With the inclusion of 2 wt % Cu-Al 2 O 3 , the thermal stability of the composite is improved, where T max raises from 397 to 405 • C and T f from 480 to 492 • C. At 14 wt % filler loading, T max and T f are respectively 30 • C and 9 • C higher compared to the neat polymer. At this point (T f 489 • C), 67% residue is still left. Upon heating the polymer, the long chains break down into small fragments which might have interacted with Cu-Al 2 O 3 particles and got trapped into filler particles difficult to be decomposed further, thus improving the thermal stability of the PS-b-(PE-r-B)-b-PS-g-MA composites [56]. Similar behavior was observed previously, where thermal stability was enhanced due to filler incorporation which hindered the segmental movement of polymer when intermingled with small chains of the host polymer [27,48]. Analogous degradation patterns are seen in the derivative thermogravimetry (DTG) curves of host polymer and its composites ( Figure 12).

Thermal Characteristics of Block Copolymer Composites
3.6.1. Thermogravimetric Analysis (TGA) TGA examines the thermal properties as the weight alteration upon heating during the phases of thermal breakdown. The thermal behavior determines the possible specific application fields of nanocomposites [59]. TGA thermograms of neat PS-b-(PE-r-B)-b-PS-g-MA and Cu-Al2O3 loaded composites from 0, 2, and 14 wt % are shown in Figure 11. A two-phase decomposition is observed for neat block copolymer. A slight dip at 250 °C indicates the presence of some residual low molecular weight compounds in the polymer. In the present conditions, the polymer remains stable up to 397 °C (8% weight loss). The second phase of decomposition starts at 397 °C (Tmax) and continues up to a final degradation temperature of 480 °C (99% weight loss at Tf). With the inclusion of 2 wt % Cu-Al2O3, the thermal stability of the composite is improved, where Tmax raises from 397 to 405 °C and Tf from 480 to 492 °C. At 14 wt % filler loading, Tmax and Tf are respectively 30 °C and 9 °C higher compared to the neat polymer. At this point (Tf 489 °C), 67% residue is still left. Upon heating the polymer, the long chains break down into small fragments which might have interacted with Cu-Al2O3 particles and got trapped into filler particles difficult to be decomposed further, thus improving the thermal stability of the PS-b-(PE-r-B)-b-PS-g-MA composites [56]. Similar behavior was observed previously, where thermal stability was enhanced due to filler incorporation which hindered the segmental movement of polymer when intermingled with small chains of the host polymer [27,48]. Analogous degradation patterns are seen in the derivative thermogravimetry (DTG) curves of host polymer and its composites ( Figure 12).  DSC analysis provides the determination of the glass transition temperature (Tg) of materials, the temperature at which a polymer transforms from a glassy to a rubbery state [60]. The DSC thermograms of neat PS-b-(PE-r-B)-b-PS-g-MA and the corresponding composites with 0, 2, and 14 wt % filler loading are shown in Figure 13. The stiffness of polymers is usually studied by Tg analysis. Stiff polymer chains with bulky, rigid side groups attached to the main chain imparts a high Tg. It is known that at Tg, polymer chains start to move. The results show that the incorporation of Cu-Al2O3 in the polymer matrix increases the Tg as the chains of PS-b-(PE-r-B)-b-PS-g-MA strongly adhere to the Cu-Al2O3 particles, which prevents free motion of polymer chains and hinders the segmental movement of chains [61,62].

Conclusions
In this study, nanocomposites were synthesized from the block copolymer polystyrene-blockpoly(ethylene-ran-butylene)-block-polystyrene-graft-maleic anhydride (PS-b-(PE-r-B)-b-PS-g-MA) as the matrix and from a filler material, prepared by the electroless deposition (ELD) of Cu particles on alumina powder. The nanocomposite belongs to the class of inorganic-organic composites containing

Differential Scanning Calorimetry (DSC)
DSC analysis provides the determination of the glass transition temperature (T g ) of materials, the temperature at which a polymer transforms from a glassy to a rubbery state [60]. The DSC thermograms of neat PS-b-(PE-r-B)-b-PS-g-MA and the corresponding composites with 0, 2, and 14 wt % filler loading are shown in Figure 13. The stiffness of polymers is usually studied by T g analysis. Stiff polymer chains with bulky, rigid side groups attached to the main chain imparts a high T g . It is known that at T g , polymer chains start to move. The results show that the incorporation of Cu-Al 2 O 3 in the polymer matrix increases the T g as the chains of PS-b-(PE-r-B)-b-PS-g-MA strongly adhere to the Cu-Al 2 O 3 particles, which prevents free motion of polymer chains and hinders the segmental movement of chains [61,62]. DSC analysis provides the determination of the glass transition temperature (Tg) of materials, the temperature at which a polymer transforms from a glassy to a rubbery state [60]. The DSC thermograms of neat PS-b-(PE-r-B)-b-PS-g-MA and the corresponding composites with 0, 2, and 14 wt % filler loading are shown in Figure 13. The stiffness of polymers is usually studied by Tg analysis. Stiff polymer chains with bulky, rigid side groups attached to the main chain imparts a high Tg. It is known that at Tg, polymer chains start to move. The results show that the incorporation of Cu-Al2O3 in the polymer matrix increases the Tg as the chains of PS-b-(PE-r-B)-b-PS-g-MA strongly adhere to the Cu-Al2O3 particles, which prevents free motion of polymer chains and hinders the segmental movement of chains [61,62].

Conclusions
In this study, nanocomposites were synthesized from the block copolymer polystyrene-blockpoly(ethylene-ran-butylene)-block-polystyrene-graft-maleic anhydride (PS-b-(PE-r-B)-b-PS-g-MA) as the matrix and from a filler material, prepared by the electroless deposition (ELD) of Cu particles on alumina powder. The nanocomposite belongs to the class of inorganic-organic composites containing

Conclusions
In this study, nanocomposites were synthesized from the block copolymer polystyrene-blockpoly(ethylene-ran-butylene)-block-polystyrene-graft-maleic anhydride (PS-b-(PE-r-B)-b-PS-g-MA) as the matrix and from a filler material, prepared by the electroless deposition (ELD) of Cu particles on alumina powder. The nanocomposite belongs to the class of inorganic-organic composites containing metal-coated ceramic reinforcement agent embedded in a thermoplastic polymer insulation, categorized as conductive polymer nanocomposites. The nanocomposites are easy to prepare, show enhanced electrical conductivity, improved thermal stability, and mechanical properties. The pronounced increment in electrical conductivity with increased filler ratio, up to 7.692 × 10 −5 S/cm in the case of 14 wt % filler loading, indicates the formation of conductive networks within the prepared composites. A good interfacial adhesion between filler and matrix permits to improve the Young's modulus and tensile strength at 14 wt % filler loading up to 159.475 MPa and 82.889 MPa, respectively. The composites also show improved thermal stability, while heat flow measurements via DSC show a higher glass transition temperature range with higher filler inclusion. XRD patterns indicate a more crystalline phase of the composites due to addition of metallic filler. SEM micrographs of the composites illustrate a uniform Cu deposition on Al 2 O 3 and its homogeneous dispersion throughout polymer matrix when using the ELD technique. These results support the potential application of the prepared composites in electronic applications that require a prolonged shelf life, both in electronic semiconductors as in microelectronic packaging, EMI-and EDS-shielding materials, antistatic coatings for electronic, flexible IT devices, and others. Depending on the requirements of the applications, these materials may be used either in coatings or for standalone components.