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3D Graphene Foam Induces Multifunctionality in Epoxy Nanocomposites by Simultaneous Improvement in Mechanical, Thermal, and Electrical Properties Leslie Embrey, Pranjal Nautiyal, Archana Loganathan, Adeyinka Idowu, Benjamin Boesl, and Arvind Agarwal ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b14078 • Publication Date (Web): 25 Oct 2017 Downloaded from http://pubs.acs.org on October 28, 2017

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ACS Applied Materials & Interfaces

3D Graphene Foam Induces Multifunctionality in Epoxy Nanocomposites by Simultaneous Improvement in Mechanical, Thermal, and Electrical Properties Leslie Embrey‡, Pranjal Nautiyal‡, Archana Loganathan, Adeyinka Idowu, Benjamin Boesl, Arvind Agarwal* Plasma Forming Laboratory Department of Mechanical and Materials Engineering Florida International University, Miami, FL 33174

KEYWORDS: Graphene Foam, Nanocomposite, Facile Synthesis, Microstructure Engineering, Thermal Transition, Mechanical Properties, Electrical Transport, Strain Sensor

ACS Paragon Plus Environment

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ABSTRACT 3D macroporous graphene foam based multifunctional epoxy composites are developed in this study. Facile dip coating and mold casting techniques are employed to engineer microstructures with tailorable thermal, mechanical and electrical properties. These processing techniques allow capillarity-induced equilibrium filling of graphene foam branches, creating epoxy/graphene interfaces with minimal separation. Addition of 2 wt. % graphene foam enhances the glass transition temperature of epoxy from 106°C to 162°C, improving the thermal stability of the polymer composite. Graphene foam aids in load-bearing, increasing the ultimate tensile strength by 12% by merely 0.13 wt. % graphene foam in an epoxy matrix. Digital image correlation (DIC) analysis revealed that the graphene foam cells restrict and confine the deformation of the polymer matrix, thereby enhancing the load-bearing capability of the composite. Addition of 0.6 wt.% graphene foam also enhances flexural strength of the pure epoxy by 10%. 3D network of graphene branches is found to suppress and deflect the cracks, arresting mechanical failure. Dynamic mechanical analysis (DMA) of the composites demonstrated their vibration damping capability, as the loss tangent (tan δ) jumps from 0.1 for the pure epoxy to 0.24 for ~ 2 wt.% graphene foam-epoxy composite. Graphene foam branches also provide seamless pathways for electron transfer, which induces electrical conductivity exceeding 450 S/m in an otherwise insulator epoxy matrix. The Epoxy-Graphene Foam composite exhibits gauge factor as high as 4.1, which is twice the typical gauge factor for the most common metals. Simultaneous improvement in thermal, mechanical and electrical properties of epoxy due to 3D graphene foam makes epoxy-graphene foam composite a promising lightweight and multifunctional material for aiding load bearing, electrical transport and motion sensing in aerospace, automotive, robotics and smart device structures.


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1. INTRODUCTION Graphene, a unique 2D carbon allotrope has attracted attention as a nanofiller candidate for engineering polymer composites due to its excellent electrical, mechanical and thermal properties.1-6 However, very high surface energies associated with a planar graphene flake/ sheet poses processing challenges. Due to intermolecular π-π interactions, graphene tends agglomerate and form clusters,7 making their uniform dispersion in the polymer difficult. Poor dispersion and agglomeration cause deterioration of properties due to non-homogeneity in the microstructure.4,8 Graphene agglomerates act as stress concentrators and can cause failure of the composite.4 Therefore, physical and chemical dispersion techniques, such as ball milling,8 centrifugation,9

sonication,10,11

melt

mixing,12

in

situ

polymerization13,14

and

functionalization15,16 are employed to restrict re-stacking of graphene flakes and enable their uniform distribution in the matrix. But these techniques have several limitations, such as damage caused due to mechanical agitation, modification of properties due to chemical reactions and introduction of undesirable defects/artifacts. Application of these techniques is also timeconsuming and complicates the composite synthesis process. To address these issues, there has been a rising interest in developing three-dimensional forms of graphene which can be introduced in the polymer matrix without any requirement of applying dispersion techniques.17-21 The 3D architectures of graphene, such as aerogels, sponges, and foams are low in density and retain excellent properties of 2D graphene.22-24 Macroporous graphene foams are particularly important for developing polymer composites, as their cellular anatomy allows the polymer resin to infiltrate, forming a composite with a uniform distribution of graphene in the matrix. An interconnected network of nodes and branches in the graphene foam provides seamless pathways for transfer of stress, electrons, and phonons. Also, 3D graphene foam has an ultra-low density


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(< 5 mg/cm3), high surface area (~850 m2g) and superhydrophobic nature. Consequently, graphene foam is a promising filler material for developing advanced multifunctional nanocomposites. Graphene foam-polymer composites have been synthesized for a wide range of potential applications, such as strain sensors,25-28 supercapacitors,29 electrochemical biosensors,30 fuel cells,31 electromagnetic interference shielding,32,33 biocompatible scaffolds,34,35 thermal interface materials,36 damping37 and acoustic backing materials.38 This study aims to harness desirable properties of graphene foam for developing epoxybased composites which exhibit simultaneous improvement in thermal, mechanical and electrical properties. It has been demonstrated that the 3D interconnected network of graphene in the epoxy matrix leads to an improvement in electrical conductivity, flexural modulus, flexural strength and fracture toughness.39-41 Jia and co-workers fabricated a graphene foam/ epoxy composite by a three-step process involving infiltration of Ni-graphene foam by epoxy resin to form prepegs, curing of the prepegs under pressure and etching away of Ni.39 Loading of 0.2 wt.% graphene foam in the polymer matrix increased the flexural modulus and strength by 53% and 38%, respectively, enhanced the glass transition temperature by 31°C, and exhibited electrical conductivity of 300 S/m. Chen et al. synthesized graphene foam/epoxy composites by dipping technique, followed by evaporation of solvents and thermal curing under vacuum.40 An electrical conductivity of 125 S/m was reported for the composite at 2.5 vol.% filler loading. In another study, a hybrid composite was fabricated by integrating 3D graphene foam in glass fiber fabric, followed by curing in a hot press.41 The composite was found to exhibit a 70% and 206% improvement in mode I and mode II interlaminar fracture energies, respectively, and a 36% improvement in interlaminar shear strength. Crack deflection by the 3D network of graphene foam was the prominent toughening mechanism. The nature of synthesis techniques adopted in


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these studies is complex and poses some challenges: (i) the application of pressure during curing39,41 can deform and damage the 3D network of graphene foam, and (ii) the etching away of the metal template after polymer infiltration39 leads to undesirable interfacial spacing and porosity in the microstructure, which diminishes the mechanical properties. In addition, vacuum curing technique40 is time-consuming and expensive. It is imperative to develop facile synthesis routes that allow effective engineering of the composite microstructure to be able to harness the benefits of graphene foam as a filler material. Herein we report facile dip coating and mold casting techniques to fabricate epoxy/graphene foam composites in room temperature conditions. Composite microstructures are engineered to augment and tailor their properties and induce multifunctionality. The effect of graphene foam filler on the thermal, mechanical and electrical transport properties of the composites are investigated. In-situ mechanical testing of the composite under an optical microscope is performed to observe the deformation mechanisms in real time. Microstructural strain evolution is studied by digital image correlation (DIC) analysis to gain insight into the mechanics of stress-transfer. The mechanical analysis performed in this study provides fundamental information about hierarchical deformation characteristics in a 3D graphene foam based polymer composite, which will pave the way for the development of novel high-strength composites with superior load bearing capability. Finally, the excellent mechanical and electrical properties of the composite are exploited to develop strain-sensors, and their electromechanical characteristics are evaluated. The overarching objective of this work is to examine the feasibility of employing 3D graphene foam as a nanofiller for inducing multifunctionality in polymeric materials using facile synthesis techniques.


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2. EXPERIMENTAL SECTION 2.1. Fabrication: Free-standing 3D graphene foam was obtained from Graphene Supermarket (Calverton, NY, USA). Aeromarine 300/21 epoxy resin was used for this study (Aeromarine Products, Inc., CA, USA). The epoxy-graphene foam composites were synthesized by dip-coating and mold-casting techniques (Figures S1 and S2). In the case of former, freestanding 3D graphene foam specimens (40 mm x 5 mm x 0.5 mm) were dipped in epoxy resin for 1, 3, 5, 7 and 9 seconds, retracted from the resin and then the resin adhering to the foam was allowed to cure for 24 hours in the air. These dipping times were found to correspond to 1.99, 1.91, 1.85, 1.64 and 1.55 weight % graphene foam, respectively. For synthesizing the composite by casting technique, the resin was poured into a 3D printed mold (40 mm x 5 mm x 3 mm) and graphene foam was laid in between the resin. 1 and 2 layer graphene foam-epoxy composites were fabricated by mold casting, which corresponds to 0.07 and 0.13 weight % graphene foam, respectively. 2.2. Characterization: Microstructure of the composite was characterized by scanning electron microscopy of the fractured surfaces, using JEOL JSM-6330F field emission SEM (Tokyo, Japan). Glass transition temperature was determined by differential scanning calorimetry (DSC) test in SDT Q600 (TA Instruments, New Castle, USA). The DSC tests were performed in an Argon environment by heating the specimens up to 250°C at the heating rate of 10°C per minute. Tensile and flexural tests were conducted in MTI SEMtester (Albany, USA) using custom-made fixtures. The real-time deformation videos of the specimens were captured by Dino-Lite AM2111 digital microscope (New Taipei City, Taiwan). Microstructural strain analysis was performed using VIC-2D Digital image correlation software (Correlated Solutions, Irmo, USA). The electrical conductivity of the composites was measured by four probe


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technique, using KEITHLEY 2401 SourceMeter (Cleveland, USA). Electromechanical characteristics were determined by stretching the composite (using MTI SEMtester, Albany, USA) and performing electrical measurements (using KEITHLEY 2401 SourceMeter, Cleveland, USA) as a function of mechanical strain. 3. RESULTS AND DISCUSSION 3.1. Microstructure engineering. Figure 1a illustrates the macroporous architecture of as-received graphene foam which has a cellular structure and a pore size of about 580 µm. To develop graphene foam based composites, it is essential for the polymer resin to infiltrate these foam cells as well as wet the foam walls. A low-viscosity polymer with superior flowability is the ideal candidate to develop such composite structures. The capability of epoxy resin to infiltrate the graphene foam was studied by a sessile drop set-up, where a single drop of the resin was placed on the foam surface, and the eventual shape evolution was observed. The initial contact made by the drop with the foam surface was characterized by a contact angle of 67° (Figure 1b). With time, the contact angle progressively decreases and eventually becomes zero in a short period of 4 seconds. This indicates that the resin rapidly infiltrates into the open cells of the foam structure. The viscosity of the resin plays a crucial role in infiltration. Superior flowability of epoxy at room temperature provides the advantage of facile room-temperature synthesis of nanocomposites. It is noteworthy that the infiltration of the foam by polymer will also depend on the foam structure, such as cell size and dimensions of the hollow foam branches. The larger pore size of the foam and lower polymer viscosity are ideal conditions for synthesizing the composites. Therefore, we propose an infiltration factor (If) defined as the ratio of these two parameters:


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ௗ

‫ܫ‬௙ ห = ఓ

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(1)

்

Where d is the cell size of graphene foam and µ is the viscosity of the resin. A larger value of this factor would correspond to favorable infiltration. The epoxy-graphene foam pair used in this study exhibited an infiltration factor of ~1 µm/cP. To elucidate the significance of this parameter, we consider the findings of our previous work on polyimide-graphene foam nanostructures,37 where it was observed that there was virtually no filling of graphene foam cells by the polyimide resin at room temperature. On the other hand, heating the resin up to 100°C resulted in well infiltrated graphene foam-polyimide nanostructures.37 Based on the available literature, it is known that the viscosity of polyimide is 1140 cP at 40°C which drops to 300 cP at 70°C. The corresponding change in the infiltration factor would be from ~0.508 (for poor infiltration) to ~1.93 µm/cP (for excellent infiltration). This suggests the value of If around 1, or greater than 1 typically signifies good infiltration.42 As the interest in graphene foam based polymer composites is increasing, a parametric quantification will facilitate predictable processing in the future. Preservation of graphene foam’s 3D node-branch anatomy is essential to harness its desirable properties for developing multifunctional nanocomposites. The fact that epoxy favorably infiltrates graphene foam was exploited for composite synthesis. A facile fabrication route was adopted, where the foam was dipped in the resin for a certain period. It was noticed that as the dipping duration was increased, the volume of polymer infiltration and adhesion on the foam also increased. This was more prominent for first 7 seconds, as the weight fraction of graphene dropped from 1.99% for 1 second dipping time to 1.64% for 7 seconds dipping time (summarized in Table 1). However, a plateau was reached after that, as an increase in dipping time from 7 to 9 seconds barely increased the polymer fraction adhering to the foam by 0.09%.


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Further dipping resulted in negligible to none weight fraction variation, suggesting an infiltration threshold. This is evidenced by comparing scanning electron micrographs (SEM) of the fractured surface of the composites fabricated by dipping foam in the resin for 1 second and 9 seconds, as shown in Figures 1c and d, respectively. A dipping duration of 1 second allows for only nominal infiltration of polymer, and the microstructure is characterized by hollow foam branches with ~11.3% porosity (Figure 1c). However, for 9 seconds dipping duration, the foam was well infiltrated by the polymer, and most of the hollow branches were also filled (Figure 1d). This resulted in a drop in porosity to ~6.3% (Table 1). The hollow foam branches are 30-50 µm in diameter. These hollow branches will be filled by the resin due to capillary action driven by Laplace pressure (approximating the branch cross-section to be circular): ∆ܲ =

ଶఊ௖௢௦ఏ

(2)

௥

Where ∆P is the pressure difference across the liquid-vapor interface, θ is the interface contact angle, γ is the surface tension of the resin and r is the radius of the hollow branch. Laplace pressure will be very high for the ultra-fine micro-branches, resulting in the polymer filling. This is a desirable phenomenon to obtain composite microstructures with the enhanced matrix-filler interfacial contact area, allowing for the superior transfer of stress, electrons, and phonons. The higher surface tension of the polymer and lower interfacial contact angle will accelerate polymer filling and infiltration. The dip coating technique is limited to fabrication of flexible and slender graphene-rich composites (Figure S3), and cannot be employed for developing bulk and robust composites due to saturation of foam/polymer resulting in low polymer content. For structural applications, the higher volume fraction of the polymer matrix is desirable. We adopted mold casting technique to


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develop bulk composites, where graphene foam strip(s) were laid in the resin contained in a mold and then allowed to cure (Figure S2). The volume fraction of graphene was varied by varying the number of layers of graphene foam strips. In this study, 1-layer and 2-layer graphene foamepoxy composites were fabricated, yielding 0.07 and 0.13 wt.% of graphene, respectively (Table 1). SEM micrograph of the polymer-rich composite microstructure (fractured surface) is shown in Figure 1e, characterized by filled cells/ pores. These composites were ultra-lightweight, characterized by densities lower than 1.2 g/cm3 (Table 1). Good interfacial adhesion at polymer/graphene interface is essential to allow for effective transfer of stresses, electrons, and phonons to obtain superior properties of the composite. During polymerization, long monomer bonds (4-5 Å) transform into short polymer bonds (1.54 Å).37 This results in the appearance of matrix-filler interfacial spacing in the microstructure of the cured polymer composite (Figure 1f). The interfacial separation at epoxy/foam branch interface was characterized in the synthesized composites based on SEM micrographs and is plotted in Figure 1g. For dip-coated specimens, the interface separation increases with the increasing polymer content (or with decreasing graphene foam fraction), due to more noticeable shrinkage associated with higher polymer resin volume adhering to the foam surface. Mold cast composites, on the other hand, exhibited a uniform interfacial spacing of 11.5 µm (within the margins of error). This is because the volume of polymer adhering to the foam walls reaches its peak after 9 seconds of dipping and the additional polymer added during mold casting is not interacting with the graphene foam surface. Hence, there is no difference in interfacial spacing on changing the polymer volume fraction in the mold-cast composite specimens. It is noteworthy that these values indicate average interfacial spacing obtained after making multiple measurements from different micrographs. Not all polymer-graphene interfaces are characterized


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by spacings and there are regions where graphene foam walls are perfectly wrapped by epoxy. This is attested by low values of porosity in the synthesized composites, ranging from ~6% to 13% (Table 1). This ensures that the transfer of stresses, electrons and phonons takes place for superior mechanical, electrical properties and thermal properties (which are explored in subsequent sections). 3.2. Thermal transition characteristics. Epoxy used in this study was found to exhibit a glass transition temperature (Tg) of 106°C. Graphene foam-based composites showed enhanced Tg (summarized in Table 2). Merely 1.99 wt.% graphene foam addition led to an impressive increase in Tg by 56°C (a 53% improvement over pure epoxy). Presence of a secondary phase in the polymer matrix usually restricts the chain mobility, which causes the enhancement in transition temperature.43 At the same time, the presence of an additional filler also limits polymer chain cross-linking and increases their free volume. This should have a counter effect and lead to a reduction of Tg. In the case of graphene foam filler, the 3D network spans the entire matrix, significantly confining the chains mobility. The 3D macroporous network provides a very high surface area, excellent dispersion/ distribution of graphene foam in the polymer matrix and overcomes the issue of agglomeration, resulting in enhanced interactions between graphene and epoxy. Therefore, the net effect is higher transition temperature of the composites. This results in superior thermal stability of the epoxy-graphene foam composite. As the graphene foam content in the composite increases, the net volume of polymer matrix confined by the 3D network of graphene increases, which results in enhanced obstruction to chain mobility. As a result, a progressive increase in Tg was recorded with increasing filler content (Table 2). Improved thermal stability makes these composites suitable for the structural body of aircraft and other applications, which might experience in-service thermal stresses.


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3.3. Mechanical properties and mechanics of stress transfer. The 3D network of graphene provides pathways for stress transfer, while at the same time eliminating concerns about the formation of clusters of graphene flakes which usually act as stress concentrators and potential failure sites in the microstructure. The slender dip coated specimens are fragile and not suitable for structural applications (as shown in Figure S3), whereas the mold-cast composites are desirable for load-bearing applications due to their rigidity and bulkiness. Therefore, the tensile and flexural investigations were performed only for bulk mold-cast composite. The moldcast composite was subjected to uniaxial tensile testing, and a 12% enhancement of ultimate tensile strength was obtained by merely 0.13 wt.% graphene foam addition to the epoxy (stressstrain plot shown in Figure 2a). The ductility of the composite was arrested (failure strain ~ 20%) as compared to pure epoxy (failure strain ~ 47%). The tensile stretching and failure of the composite were recorded under an optical microscope to observe the deformation behavior in real time (see videos in Supplementary information, Videos S1, and S2). A blue dye was added to epoxy for enhanced visualization of the foam network under LED illumination during tensile deformation. Digital image correlation (DIC) analysis of the in-situ optical video was performed to develop insight into strengthening mechanisms associated with the graphene foam (Figures 2b-c). The resultant strain contours revealed non-uniform deformation due to the cellular architecture of graphene foam. Strain contour maps corresponding to the elastic, plastic and failure regions of the tensile stress-strain curve for the composite are shown in Figures 2b, 2c, and 2d, respectively. Each of these regimes exhibits an order of magnitude variation in the strain values. The initiation of crack was captured, and corresponding localized failure strain was quantified (20.2%), encircled in the contour map of the failure regime shown in Figure 2d. As noted above, the localized deformation was heterogeneous and can be correlated to the


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microstructural features. For instance, the regions in the composite microstructure consisting of broken cellular branches exhibited higher local strains (Figure 3a and 3b). It was also noted from the strain mapping in Figure 3b that the periphery of the contours coincided with the foam branches and hence, revealed lower strains. The inner region of the cells was characterized by higher strains due to the higher ductility of the polymer. The nature of strain distribution shows the interconnected graphene branches provide seamless pathways for stress transfer, restrict and confine the matrix deformation, and improve the overall load bearing capability of the composite. The role of localized strains in crack propagation was also studied (Figure 4). The sharp nose of strain contour predicted the crack propagation pathways (Figure 4a), as shown and compared with the optical image of the corresponding region of the failed composite (Figure 4b). There is an excessively high strain in the microstructure along the crack, explaining the role of stress-concentration at the crack tip in the failure of the composite. Examination of strain contours in conjunction with stress-strain curve provides insight into microstructural deformation mechanisms, which is critical for engineering composites with superior strength. The failure characteristics were examined by post-failure electron microscopy of the fractured surface. Observation of broken branches signifies that graphene foam bears the tensile load (Figure 5). The high magnification micrograph shows the excellent interface between graphene foam branch and epoxy filled inside it (Figure 5a), which enables stress-transfer from polymer to graphene. The stretching of the foam branch was also observed, suggesting its role in resisting failure of the composite (Figure 5b). Flexural characteristic of the composite was investigated by three-point bending to examine the fracture-resistance of the material. Due to its continuous 3D network, graphene foam reinforcement is expected to hinder crack propagation and enhance the strength. Merely 0.6


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wt.% graphene foam filler resulted in 10% improvement in flexural strength (Figure 6a). 3D graphene network deflects cracks and absorbs mechanical energy to induce strengthening. This was seen in the SEM micrograph of the failed fracture surface, where graphene foam inhibited and deflected the crack propagation during deformation (Figure 6b). Also, bending of foam branch was observed in Figure 6c, indicating energy absorption by graphene network. A very high magnification view of a foam branch shows ripples/ corrugated surface (Figure 6d). These corrugations could be intrinsic to the graphene structure, or they could be formed during deformation.44-46 These ripples/ micro-textures aid in the dissipation of mechanical energy,46-48 thereby enhancing the composite’s ability to resist fracture. 3D graphene foam is characterized by very high porosity approaching ~99% and ultralow density (less than 5 mg/cm3). As a result, graphene foam based nanocomposites have very high specific tensile strength and Young’s modulus. This makes these composites highly capable of absorbing mechanical energy.17 Therefore, epoxy-graphene foam composites should protect against localized impact. The ability of material to dissipate mechanical energy is quantified by loss tangent, which is the ratio of loss modulus (E”) to storage modulus (E’):49 ா"

tan ߜ = ாᇱ

(3)

Higher values of tan δ signify superior energy dissipation capability.50 To study energy absorption capacity of the epoxy-graphene foam composites, impact-based damping experiments were performed.37,49 A steel stylus was dropped on the specimen surface, and the bouncing probe was then naturally allowed to come to rest. These tests were conducted at 1 mN initial stylus load. A low load of 1 mN was adopted to prevent any damage to pure graphene foam specimen without the polymer backbone. To be able to capture the response from the foam and develop an


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understanding of interfacial dynamic stress-transfer phenomena, the tests were restricted to slender dip-coated specimens. The technique adopted primarily captures the mechanical response from/ around the sample surface, and a mold-cast composite with polymer-rich matrix would not reveal response of graphene foam at a very low load of 1 mN. Hence these studies were not performed on mold cast specimens. Pure epoxy had a loss tangent of ~0.1, whereas virgin graphene foam was characterized by an extraordinary tan δ approaching 0.6 (Figure 7). Dynamic loading of graphene is known to induce ripple formation and propagation.37,46 also, the probe also suppresses intrinsic corrugations of graphene, causing absorption of energy. Each foam branch is composed of multiple layers of graphene deposited during CVD. The impacting stylus compresses these layers, causing a reduction in inter-layer spacing. As a response, the inter-layer van der Waals forces exert an opposing repulsive force. This spring-like action leads to dissipation of energy, resulting in excellent tan δ values observed for graphene foam. Addition of graphene foam was found to enhance the damping potential of the epoxy composite, as tan δ increased by 140% up to 0.24 at 1.99 wt. % graphene foam reinforcement. In addition to the localized energy dissipation mechanisms associated with 2D graphene, the 3D architecture of foam will also lead to energy absorption due to bending of the branches under the influence of load. Higher filler content in the matrix leads to enhanced overall area of interaction between graphene and epoxy. As a result, stress transfer from the matrix to highly stiff graphene phase will be more effective, leading to superior mechanical properties. Moreover, due to its 3D anatomy, there is no concern related to agglomeration for higher filler fraction which is a major drawback for composites based on 2D graphene flakes. The hierarchical structure of graphene foam provides colossal opportunity to tailor the mechanics of these nanocomposites. The composites with energy


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dissipation capability are highly desirable for aerospace and automotive structures, precision systems, robotic structures, and NEMS/MEMS devices. 3.4. Electrical transport. Graphene is well known for its extraordinary electrical transport characteristics, attracting attention as a secondary filler in polymer matrices to impart electrical conductivity. However, random distribution of the graphene flakes in the polymer restricts the capability to allow electron mobility. A 3D network of graphene is expected to allow seamless pathways for transfer of electrons in the matrix. Epoxy used in this study exhibited a very low conductivity of the order of 10-9 S/m. The graphene foam was characterized by a conductivity of 1450 S/m. Graphene foam reinforcement resulted in electrically conductive composites, with tailorable conductivities as shown in Figure 8. As the graphene foam content in the matrix increases, there is an effective increase in the interconnected pathways for the seamless transfer of electrons to take place. This results in improved electrical conductivity for higher filler fraction. The conductivity as high as 461 S/m was obtained for 1.99 wt.% graphene foam addition, which is 11 orders of magnitude higher than pure epoxy! The ability to enhance electrical transport remarkably at such low nanofiller loading is highly desirable for application in sensors, batteries, electronic devices and stretchable electronics platforms. 3.5. Electromechanical characteristics. Excellent mechanical stability coupled with brilliant electrical transport properties of these composites paves the way to motion-sensing applications of the material. This is accomplished by exploiting the variation in electrical resistance of the material due to mechanical deformation. There has been considerable interest in recent years in the area of graphene foam-polymer composites based strain sensors for body motion sensing.25-28 Flexible composites based on PDMS have been developed for stretchable electronics.25-28,51 However, graphene foam-composites also have potential application as rigid


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and robust sensors for ‘structural’ motion-sensing. Epoxy, being a structural polymer, is an ideal candidate for fabricating graphene foam based mechanically strong and electrically conductive strain sensors. The electromechanical behavior of graphene foam-epoxy composite was examined by subjecting it to tensile stretching and measuring the electrical conductivity of the strained composite. Figure 9 shows the variation of the relative resistance change with respect to the applied mechanical strain. The electromechanical sensitivity was probed in low-to-medium range mechanical strains (up to ~20%). These tests were conducted for both dip-coated (1.9 wt.% graphene foam) and mold-cast (0.1 wt.% graphene foam) composites. The increase in resistance with applied mechanical strain was noticed, which is attributed to stretching of graphene foam cells followed by failure of the foam branches, disrupting conduction pathways. Gauge factor, which is a measure of the sensitivity, was evaluated as: GF =

∆ୖ/ୖ౥

(4)

க

Where Ro is the resistance at 0% strain, ∆R is the change in resistance with respect to Ro and ε is the applied strain. The gauge factor was computed by determining the slope of resistance variation versus strain plots. Gauge factor values were obtained as 2.6 and 4.1 for dip coated and mold cast composites, respectively. As discussed in Section 3.1, filling of cells and voids in graphene foam during mold casting leads to a rigid polymer-rich microstructure. Therefore, during tensile deformation, the graphene foam branches are not easily damaged, which prevents the dramatic increase in the resistance of the composite. Contrary to this, slender dip-coated specimens are highly flexible (Figure S3), and not as rigid as the mold-cast samples. Therefore, their resistance-retention capability is inferior, resulting in a lower gauge factor. Nevertheless, the values for both the composites exceed the typical Gauge factor value of 2.0 for the most


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commonly used metallic alloys of Cu, Ni, Cr, Fe and Al in strain gauges.52 Therefore, graphene foam based lightweight polymeric composites are ideal candidates for replacing bulkier metallic strain gauges, having positive implications for smart and lightweight aerospace, automotive and robotic structures. 4. CONCLUSIONS In this study, 3D graphene foam based multifunctional epoxy composites were fabricated by dip coating and mold casting techniques. A new processing parameter was “infiltration factor” is introduced, which relates polymer viscosity and 3D graphene foam geometry to predict the processability of the composites. An infiltration factor around or greater than unity signifies active infiltration of foam cells by a polymer resin. •

Addition of 2 wt. % graphene foam was found to increase the glass transition temperature of epoxy from 106°C to 162°C, improving the thermal stability of the polymer composite.

•

A small volume fraction of Graphene foam addition resulted in improvement in the mechanical properties including e ultimate tensile strength, flexural strength, and loss tangent. 3D network of graphene branches suppress and deflects cracks and restricts deformation.

•

In situ mechanical investigation coupled with digital image correlation (DIC) analysis provided an understanding of stress transfer behavior due to cellular reinforcement. Strain contours in DIC maps indicate that graphene foam cells restrict and confine the deformation of the matrix, thereby enhancing the load-bearing capability of the composite.


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•

Graphene foam branches also provide seamless pathways for electron transfer, which induces a remarkable electrical conductivity exceeding 450 S/m (for 2 wt. % graphene foam addition) in an otherwise insulator epoxy matrix.

•

The composite exhibited impressive electromechanical characteristics, with gauge factor as high as 4.1, which is double than the typical gauge factor for metallic strain gauges.

Macroporous graphene foam enables the development of hierarchical nanocomposite microstructures with tailorable thermal, mechanical, electrical and electromechanical characteristics.


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FIGURES


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Figure 1. (a) SEM micrograph of as-received macroporous 3D graphene foam fabricated by chemical vapor deposition. (b) Epoxy resin/graphene foam interfacial contact observed optically in a sessile single drop experiment. SEM micrographs of the fractured surface of dip-coated epoxy/graphene foam composite fabricated for dipping times of: (c) 1 second and (d) 9 second, showing hollow and filled graphene foam branches, respectively. (e) SEM micrograph of the fractured surface of a polymer-rich mold-cast microstructure comprising of fully filled foam cells/pores. (f) High-resolution FESEM image is showing graphene foam/epoxy interfacial spacing. (g) Quantification of average interfacial spacing in the composites with variable graphene foam weight fractions fabricated by mold-casting and dip-coating techniques.


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Figure 2. (a) Tensile stress-strain curve for pure epoxy and mold-cast epoxy/graphene foam composite. Strain contour maps corresponding to: (b) elastic, (c) plastic and (d) failure regions of the stress-strain curve.


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Figure 3. Correlation of mechanical deformation with microstructure feature established by comparing instantaneous optical captured during in-situ testing (a) with the corresponding strain contour map obtained by digital image correlation analysis (b).


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Figure 4. Nose of the strain contour map during failure predicts the crack propagation pathway (a), attested by the final optical image of the failed composite specimen (b).


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Figure 5. Examination of failure characteristics by SEM reveals: (a) strongly adhered polymergraphene interface post failure evidencing effective stress-transfer, and (b) stretching of graphene foam due to tensile loading, showing the effectiveness of flexible foam in resisting failure.


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Figure 6. (a) Flexural stress-strain plot showing graphene foam induced strengthening. (b) SEM micrograph of composite cross-section post-failure showing suppression and deflection of cracks due to graphene foam. (c) SEM micrograph showing signs of interfacial stress-transfer and bending of foam branch, which evidences the role of graphene foam in absorbing mechanical energy and enhancing the strength. (d) High magnification FESEM image of graphene foam branch showing ripples/ corrugations, which assist in energy dissipation due to interface friction.


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Figure7. Loss tangent values for dip-coated epoxy-graphene foam composites, as a function of graphene foam weight fraction. Loss tangent values for pure epoxy and graphene foam are also plotted.


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Figure 8. Electrical conductivity values for epoxy-graphene foam composites fabricated by dipcoating as well as mold-casting, as a function of graphene foam weight fraction. Conductivity value for graphene foam is also plotted.


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Figure 9. Relative change in resistance of epoxy-graphene foam composite as a function of tensile strain for: (a) dip-coated composite (1.9 wt. % graphene foam) in low strain regime, and (b) mold-cast composite (0.1 wt. % graphene foam) in medium strain regime.


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TABLES Table 1. Processing approaches for developing Epoxy-Graphene foam composites Synthesis Approach

Graphene foam wt. fraction (%)

Mass density (g/cm3)

Porosity (%)

1 second

1.99

1.015

11.303

3 seconds

1.91

1.021

10.658

5 seconds

1.85

1.026

10.189

7 seconds

1.64

1.042

8.676

9 seconds

1.55

1.068

6.347

1-layer composite

0.07

1.135

12.96

2-layers composite

0.13

1.134

13.35

Processing Parameter Dipping duration

Dip coating

Graphene foam layers Mold casting


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Table 2. Glass transition temperature of graphene foam-epoxy composites Graphene foam wt. fraction (%)

Tg (in °C)

0

106

0.07

110

0.13

118

1.55

149

1.64

152

1.85

155

1.91

159

1.99

162

ASSOCIATED CONTENT Supporting Information. Supporting Figures, S1-S3. Schematic representation of composite synthesis techniques and illustration of the flexibility of slender dip-coated composite. Supporting Videos S1 and S2. Tensile stretching and failure of Epoxy/Graphene foam composite.


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AUTHOR INFORMATION Corresponding Author *Arvind Agarwal Plasma Forming Laboratory Department of Mechanical and Materials Engineering Florida International University Miami, FL, USA 33174 E-mail: [email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have approved the final version of the manuscript. ‡These authors contributed equally. Funding Sources W911NF-15-1-0458 grant by the Air Force Office of Scientific Research.

ACKNOWLEDGMENT The authors would like to acknowledge US Department of Defense W911NF-15-1-0458 grant. Pranjal Nautiyal thanks, Florida International University Graduate School for the financial support through Presidential Fellowship. Advanced Materials Engineering Research Institute (AMERI) at Florida International University is acknowledged for the research facilities used for this study.


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