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Applications of Polymer, Composite, and Coating Materials

Roll-to-Roll Processing of Silicon Carbide Nanoparticle Deposited Carbon Fiber for Multifunctional Composites Christopher C. Bowland, Ngoc A Nguyen, and Amit K Naskar ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b03401 • Publication Date (Web): 13 Jul 2018 Downloaded from http://pubs.acs.org on July 17, 2018

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ACS Applied Materials & Interfaces

Roll-to-Roll Processing of Silicon Carbide Nanoparticle Deposited Carbon Fiber for Multifunctional Composites Christopher C. Bowland*, Ngoc A. Nguyen, Amit K. Naskar* Carbon and Composites Group, Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37931, USA Keywords: Silicon carbide, nanoparticles, carbon fiber, multifunctional composites, structural health monitoring Abstract This work provides a proof of principle that a high volume, continuous through-put fiber coating process can be used to integrate semiconducting nanoparticles on carbon fiber surfaces to create multifunctional composites. By embedding silicon carbide nanoparticles in the fiber sizing, subsequent composite fabrication methods are used to create unidirectional fiber reinforced composites with enhanced structural health monitoring (SHM) sensitivity and increased interlaminar strength. Additional investigations into the mechanical damping behavior of these functional composites reveals significantly increased loss factor at the glass transition temperature ranging from a 65% to 257% increase. Composites with both increased interlaminar strength and SHM sensitivity are produced from a variety of epoxy and silicon carbide nanoparticle concentrations. Overall, the best performing composite in terms of combined

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performance shows an increase of 47.5% in SHM sensitivity and 7.7% increase in interlaminar strength. This work demonstrates successful and efficient integration of nanoparticle synthesis into large-scale, structural applications.

1. Introduction Nanomaterials have garnered much of the research spotlight in recent years by promising profound improvements in material performance across a wide variety of industries. This hype was initially sparked by the synthesis of carbon nanotubes (CNTs),1, 2 and further incited after the first isolation of single layer graphene.3 Specifically, the outstanding tensile strength of CNTs led to the moonshot idea of constructing a space elevator that resonated with the public and caused the nanomaterial hype to explode.4 Despite the vast amount of research devoted to these carbon nanomaterials and their ever-increasing inclusion in patents, commercialization of these highly-praised materials have not lived up to all of the expectations mainly due to high costs related to synthesis challenges in controlling their dimensions and chirality, difficulties with large-scale integration and possible health concerns from the nanomaterials.5,

6

However,

research has made great strides in high-level control over the architecture, dimension and surface chemistry of various ceramic nanomaterials using more cost-effective routes, which now offer new alternatives to CNTs and graphene. To name a few, these nanomaterial manufacturing routes include electrospinning,7 hydrothermal/solvothermal processing,8 chemical vapor deposition9 and methods involving colloidal,10 template-assisted,11, 12 plasma-assisted,13 thermal decomposition14, 15 and polymer brush assisted syntheses.16, 17 Additional computational methods recently revealed hundreds of materials that can theoretically exist as 2D nanomaterials thus further expanding nanomaterial experimentation.18 These various approaches exemplify the sheer

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amount of research effort focused on developing synthesis techniques for nanomaterials with unique and exceptional properties. The challenge now lies in effectively integrating these nanomaterials into useful structures to offer scalability and economic viability to progress towards commercialization. One attractive integration approach is to utilize fibers as the nanomaterial growth substrates for use in fiber reinforced composites. Numerous ceramic nanomaterials including silicon carbide,19 lead zirconate titanate,20 barium titanate21-23 and zinc oxide24, 25 have been synthesized on fiber surfaces to contribute to increased interfacial strength and/or to add electromechanical functionalities. However, most of these synthesis protocols have harsh parameters and may require batch processing that limits fiber compatibility and severely limits fiber processing quantities.22 Additionally, these synthesis techniques would require significant efforts to optimize the scale-up process, and the final process would be applicable to only a select material system. The lab-scale processes have created remarkable nanostructures on fibers but are not feasible for high-volume, continuous fiber production, which is needed for commercialization of nanomaterial enhanced fiber reinforced composites. Some research efforts have even made it a goal to use ceramic nanowires or CNTs as a replacement for polymer fiber sizing.26, 27 Polymer sizing is applied to tows of fibers to increase the adhesion between the fiber and matrix, which CNTs and nanowires have accomplished without requiring a polymer sizing.28 However, a secondary purpose of sizing is to enhance the ease of fiber handling during subsequent fabric weaving and composite layup. Exposed ceramic nanowires and CNTs on the fiber surface only complicate downstream fiber processing by increasing the difficulty of fiber handling and increasing the risk of damaging the topology of the nanostructures. Therefore, practical solutions will rely on a polymer-based fiber phase. In 3 ACS Paragon Plus Environment


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commercial carbon fiber manufacturing, the sizing processing is one of the least costly in terms of both time and energy expenditures so it is detrimental to add complicated and time-intensive processing procedures, i.e. CNT and nanowire growth, at that manufacturing step.29 Therefore, simple modifications to the existing polymer is the appropriate approach. Previous research has demonstrated dispersing CNTs on glass fibers via a polymer sizing 3032

and dispersing them on sized carbon fibers using water.

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That previous research revealed

increased interfacial shear strength and structural health monitoring (SHM) capabilities.30-32, 34 However, little to no large-scale commercial applications have been realized thus far due to both the current high cost and health concerns of CNTs.6 In additional research efforts, silicon dioxide nanoparticles (NPs) and graphite nanoplatelets were adhered to carbon fiber via an ethanol dispersion method to increase exclusively the interfacial shear strength of the fibers,.35, 36 While that past research showed promising results for glass fiber composites, there has been no attempt to integrate silicon carbide (SiC) NPs into carbon fiber sizing. Due to the semiconducting nature of SiC, it has inherent piezoresistive behavior, which was utilized for this research.37, 38 The key to bridging this gap between nanomaterial synthesis and composite integration is to make the nanomaterial and fiber synthesis procedures independent while creating a quick, simple, and versatile integration approach. The research here demonstrates the use of commercially-available SiC NPs in a widely-used epoxy sizing to coat carbon fiber tows in a continuous throughput process. It connects high performance carbon fiber composites with lower cost ceramic NPs without the need for chemically modified surfaces to create both mechanical and SHM enhancements. Additionally, using water and only small quantities of epoxy avoids toxic solvents making this an environmentally friendly process further increasing the ease of adoption by industries and other 4 ACS Paragon Plus Environment

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researchers. This will help spark future research in integrating various nanomaterials into fiber reinforced composites, and not just carbon fiber composites. Additionally, this process promotes commercial adoption by demonstrating a highly-scalable technique that is compatible with existing continuous fiber processing lines. Consequently, nanomaterial research efforts now have an additional platform for commercial deployment, thus finally giving these novel materials more means to be applied in real world, commercial structural applications. 2. Results and Discussion A straight-forward integration procedure was used to adhere SiC NPs to the carbon fiber surface via a continuous feedthrough sizing process. As illustrated in Figure 1, a commercial epoxy sizing was obtained and diluted with deionized (DI) water to decrease the fiber coating thickness. The epoxy was supplied as an aqueous dispersion so increasing the water content did not adversely affect the epoxy chemistry. As shown in the following sections, the aqueous epoxy solution was diluted at two concentrations, 1:10 and 1:40 of epoxy to water calculated by weight. This not only helped with varying the sizing thickness but also demonstrated the versatility of this technique to accommodate different sizing viscosities and NP concentrations. After dilution, SiC NPs (Figure 1b) were weighed out as a percentage of the dehydrated epoxy and added directly to the aqueous epoxy solution. This mixture was stirred and sonicated until welldispersed. The NP concentration ranged from 0 to 50 wt.% in 10 wt.% intervals for both epoxy dilutions. This weight fraction takes into account only the dehydrated epoxy and the NPs, so this is a minuscule NP concentration when compared to the overall fiber weight as shown in the thermogravimetric analysis (TGA) in Supporting Information Figure S1. In fact, even at the highest SiC concentration tested in this paper, the SiC NP concentration is less than 3 wt.% of the entire weight of the fiber. This epoxy mixture was then added to the fiber surface through a 5 ACS Paragon Plus Environment


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simple dip coating procedure where the carbon fiber tow was submerged in the liquid then removed and dried at about 120°C in air. Scanning electron microscopy (SEM) images of the asreceived bare carbon fiber and the resulting SiC NP deposited fibers can be seen in Figure 1a,c, respectively. Upon drying, the tow was collected on a spool to await composite fabrication. This was all performed in a continuous feedthrough process that successfully maintained ease of fiber handling in the subsequent composite fabrication steps despite the inclusion of NPs on the fiber surface. The coated fibers were then used to make unidirectional fiber reinforced composites with 60 vol.% fibers via compression molding using an epoxy matrix material. This resulted in 12 different composite compositions that were sectioned with various dimensions for SHM, mechanical and viscoelastic property characterization. More detailed descriptions of the fiber coating and composite fabrication processes be found in the Experimental Procedures section.

Figure 1. Schematic showing the NP coating process. The epoxy emulsion is initially diluted with water then the SiC NPs are added and mixed to get a well-dispersed mixture. Bare carbon fiber tows are then dip coated in the mixture and dried. SEM images show the (a) bare carbon fiber tow before dip coating, (b) SiC NPs and (c) SiC NP coated tow of fibers. 6 ACS Paragon Plus Environment

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2.1 Structural Health Monitoring The main motivation for using SiC NPs was to enhance the composite’s SHM capabilities. The inherent piezoresistive attribute of SiC was utilized to enhance the change in electrical resistance of the composite under varying external stresses. Due to the use of the piezoresistive effect instead of the piezoelectric effect, static loading conditions could be utilized as opposed to dynamic loads. Piezoelectric materials have been used for SHM in the literature but they require dynamic loading cycles in order to generate a signal so they cannot be used for static loading scenarios.23, 25 This enables piezoresistive SHM sensors to be employed in different real-world sensing applications than piezoelectric sensors. One condition for utilizing this piezoresistive type SHM is that the bulk material must have sufficiently low electrical resistance. This criterion makes carbon fiber a perfect candidate due to its electrically conductive property that directly translates to an electrically conductive composite. To characterize this SHM capability, the unidirectional composites were sectioned into cantilevers and devised electrodes to measure the out-of-plane through-thickness electrical resistance as illustrated in Supporting Information Figure S2a. These composites were placed in a dynamic mechanical analyzer (DMA) in a single cantilever configuration to apply repeated cycles of static strains at various magnitudes. Monitoring the out-of-plane through-thickness electrical resistance revealed clear changes in resistance as a function of applied stress. Typical strain cycles and relative electrical resistance changes are shown in Figure 2a. These loading cycles were used to calculate an average relative resistance change as a function of the strain as shown in Figure 2b,c. For each strain value, the relative resistance change was averaged over at least 10 strain cycles, and this was repeated for at least 10 different strain levels. These 100 strain



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cycles were completed for all 12 composites with varying concentrations of NPs with the 1:10 and 1:40 epoxy dilutions plotted in Figure 2b,c, respectively. For the 1:10 epoxy dilution fibers (Figure 2b), the average electrical resistance change increased significantly at a concentration of 20 wt.% (in 1:10 epoxy dilution) SiC NPs. The other concentrations either stayed roughly the same as the epoxy only case or experienced a slight decrease in electrical response magnitude. A similar trend was seen for the 1:40 epoxy dilution composites in Figure 2c. For these composites, the maximum increase in SHM sensitivity was seen in the 40 wt.% SiC NPs. This showed that the peak performance level shifted to a higher SiC NP concentration when the dilution ratio increased. As shown in the Supporting Information Figure S1, the weight fraction of epoxy on the fiber decreases as the epoxy dilution increases from 1:10 to 1:40, as expected. This repeated trend for both dilution levels confirmed that there is an optimum SiC NP concentration. Depending on the fiber sizing concentration, the optimum NP concentration will vary, at least for the case of the SiC NPs in this research using the described application technique. While the findings were not the expected results, this still illustrates that SiC NPs at the appropriate concentrations can bring about significant enhancements in the SHM capabilities of carbon fiber composites. This trend did not match with the original hypothesis that increasing the SiC concentration would further enhance the sensing sensitivity. Rather, the sensitivity reached an optimum value at specific SiC concentrations depending on the epoxy concentration. In fact, excessive amounts of SiC NPs proved to have a negative effect on the SHM sensitivity. A similar trend was reported previously in the literature for SHM with CNTs in that increased loading of CNTs in a matrix does not necessarily increase the electrical resistance change sensitivity.39-42 There was an optimum CNT loading concentration that was attributed to percolation effects. In this work, the optimum level of SiC 8 ACS Paragon Plus Environment

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loading is attributed to NP agglomerations forming on the fibers instead of strictly percolation effects. When mixing the NPs in solution, the NPs could agglomerate and settle out of the solution at higher concentrations or these agglomerations could end up on the fibers as seen in Supporting Information Figure S3. This agglomeration effect is further discussed in the following sections regarding the composite’s mechanical performance. The two relative resistance change versus strain plots (Figure 2b,c) showed a general trend in the data, but a more quantified method was used to effectively summarize the results. The typical way of characterizing sensor sensitivity is to calculate the gauge factor, which takes the relative resistance change and divides it by the applied strain as shown in Eq 1

=

(1)

where ∆R is the change in resistance, R0 is the initial resistance and ε is the strain. Following Equation 1, the overall average gauge factor for each composite was calculated. The gauge factor at each strain value was calculated using Equation 1 then all of the gauge factors over the entire strain range (0.01% - 0.12%) were averaged resulting in a single gauge factor value for each composite. This helped to better compare the sensitivity of each composite and is plotted in Figure 2d with the two epoxy dilutions denoted in different colors. The tabulated values are shown in the Supporting Information Table S1 and Table S2. Overall, the maximum gauge factors for the composites were 5.70 for the 1:10 epoxy dilution with 20 wt.% SiC NPs and 4.89 for the 1:40 epoxy dilution with 40 wt.% SiC NPs. For comparison, in the literature fiber sensors have been fabricated using CNTs and revealed a gauge factor of 1.6, which is similar to a gauge factor of about 2 for foil strain gauges, while SiC itself has shown a gauge factor of 25 to 28.38, 43 When compared to the composites with no NPs, this equated to a maximum SHM sensitivity 9 ACS Paragon Plus Environment


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improvement of 53.6% and 47.5% for the 1:10 and 1:40 epoxy dilutions, respectively. Ultimately, this showed that SiC NPs embedded in epoxy sizing and coated on carbon fiber can significantly enhanced the SHM sensitivity of the composites.

Figure 2. (a) The typical electrical resistance response to repeated input strain. Plots of the average relative resistance change as a function of input strain for various SiC NP concentrations in the (b) 1:10 dilution and (c) 1:40 dilution. (d) The average gauge factor for composite over the entire strain range tested. Error bars were calculated as a standard deviation for (b) and (c) and a 95% confidence interval for (d). 2.2 Mechanical Performance To be classified as a truly multifunctional composite, it must maintain its structural integrity while having added functionalities. Ideally, both the mechanical performance and functional 10 ACS Paragon Plus Environment

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properties will be enhanced. Since increases in the functional properties were shown in the previous section, the composite’s mechanical performance was investigated to show that the NPs created a truly multifunctional structural composite. 2.2.1 Short Beam Strength An important property of fiber reinforced composites is the interlaminar strength. While the fibers are very strong under tensile loads parallel to the fiber axis, composite weaknesses come from the fiber-to-matrix adhesion thus illustrating the need to characterize the interlaminar strength through short beam shear tests. This test uses a three-point bend configuration with a short span length to determine the composite’s interlaminar strength. Figure 3a,b show the results of these tests for the 1:10 and 1:40 epoxy dilution composites, respectively. Representative stress-strain plots in Figure 3c illustrate the characteristic curve shapes from these tests. It is seen that for the 1:10 epoxy dilution (Figure 3a) three composites experienced slight decreases in strength, but these decreases were at most a 4.4% drop in strength. More importantly, the 40 wt.% NP composite actually saw a 5.1% increase in short beam strength while the 30 wt.% NP composite retained its mechanical strength. The 1:40 epoxy dilution composites revealed even more profound strength enhancements. In fact, all of the 1:40 composites (Figure 3b) showed increases in short beam strength with the highest increase being 13.7% for the 20 wt.% NP composite. The tabulated values for these tests can be seen in the Supporting Information Table S3 and S4 with accompanying schematics showing the sectioning of the composites and test setup for these tests in Supporting Information Figure S2b,c.



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These interlaminar strength enhancements were surprising due to the lack of chemical interaction expected between the NPs and the fiber and the matrix. As a result of the NPs not having any surface functional groups, the variations in the mechanical performance can be attributed to physical contributions, such as mechanical interlocking, crack pinning, crack deflection and nanovoid formation.44-47 Mechanical interlocking has been explored for carbon fiber through etching to increase surface roughness, but these processes are too time-intensive and costly for large-scale production processes and can degrade the tensile strength of the fibers.48 As stated in the previous section, higher concentrations of NPs in the sizing mixture caused some agglomeration of the NPs. While these could serve as stress concentration sites for cracks to form, the literature has reported mechanical benefits of having this type of NP clustering in nanocomposites by creating crack pinning or crack deflection mechanisms.44, 49-51 This mechanism can change the crack propagation direction and lead to increased fracture strength. Representative SEM images of the composites’ fracture surface are shown in Supporting Information Figure S4. In those images, it is clearly seen that the composite without nanoparticles experiences complete removal of the epoxy matrix from the fiber surface in some locations leaving behind smooth fiber surfaces. However, the composites with nanoparticles always show a rough fracture surface that is likely caused by direction changes in crack propagation caused by the pinning effect of SiC nanoparticles.50 Therefore, agglomerations were detrimental to the SHM sensitivity, yet they proved to be beneficial for mechanical strength.


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Figure 3. Short beam strength tests for (a) 1:10 epoxy dilution composites and (b) 1:40 epoxy dilution composites. (c) Representative stress-strain curves from the short beam shear tests for the 1:40 epoxy dilution composite without nanoparticles and the 1:40 epoxy dilution composite with 20 wt % SiC NPs. Strain was calculated as the crosshead displacement over the sample thickness. Error bars are shown as a 95% confidence interval. This mechanical interlocking effect was validated by confirming that no additional chemical interactions were present in the composites after SiC addition. The chemical bonds present on the fibers and in the composites were identified via Fourier-transform infrared spectroscopy (FTIR). Figure 4a shows the spectra for fiber tows before and after epoxy sizing treatments while Figure 4b shows the same fiber tows after being embedded in the epoxy matrix to form



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the final composite. The fibers selected for this analysis were the as-received bare carbon fiber, the epoxy only fibers for both dilutions, the 20 wt.% in the 1:10 dilution and the 40 wt.% in the 1:40 dilution. These fibers were sufficient to provide a representative sampling of the entire fiber set. There appears to be no appreciable difference between the fibers with and without the embedded SiC NPs. As compared to the bare carbon fiber, the coated fibers show slight epoxy peaks from the sizing. These peaks were low intensity due to the low volume of sizing on the fibers, but they are slightly more intense in the 1:10 epoxy dilution fibers resulting from the increased sizing thickness. The epoxy peaks are much more prevalent in Figure 4b due to the large increase in epoxy volume. Overall, no additional chemical bonding was observed from integrating the SiC NPs thus supporting the claim that the increase in mechanical performance is due to physical contributions, in terms of mechanisms that hinder crack growth and propagation, and not chemical interactions. Therefore, the results presented here offer a simple and straight forward route to introduce mechanical interlocking to increase the interlaminar strength of carbon fiber reinforced composites via NP-embedded sizing.

Figure 4. FTIR spectra for fibers with different sizing compositions (a) before embedding in epoxy matrix and (b) after composite fabrication. 14 ACS Paragon Plus Environment

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2.2.2 Damping In addition to interlaminar strength enhancements, the viscoelastic properties of the composites showed large variations resulting from SiC NP integration. The viscoelastic properties were investigated via DMA testing using a three-point bend configuration. This is a common test to measure the storage modulus and loss modulus of polymers and composites. The storage modulus provides information regarding the elastic behavior of the material while the loss modulus is a measure of the viscous properties. Thus, this test provides the viscoelastic properties of a material. Determining those two moduli, a third value called the damping loss factor (tan δ) can be calculated by dividing the loss modulus by the storage modulus. This value provides the energy dissipation capability of a material. A higher tan δ value represents higher damping abilities, which means the material can absorb more energy. Higher damping is a sought-after material characteristic in automobiles and aircraft for vibration control and fatigue mitigation.49 Consequently, this research concentrated on comparing the tan δ of the composites with the aim of increasing that value. The composite beams were subjected to a 10 Hz sinusoidal input waveform with a strain amplitude of 0.01% and heated from room temperature to 250 °C at a rate of 2 °C/min. The resulting tan δ plots for the 12 different composites are shown in Figure 5 with plots of magnified views before (Figure 5c,f) and around (Figure 5b,e) the onset of the glass transition temperature. Around the glass transition temperature, a clear peak in tan δ is experienced as expected for a polymer as it transitions from the glassy state to a rubbery state. Results showed increased damping at the glass transition temperature for the NP integrated composites. Both the 1:10 and 1:40 dilution composites showed significant increases in the tan δ at the glass transition as shown in Figure 5b,e. For the 1:10 composites (Figure 5a,b), the peak tan δ value was 0.13 15 ACS Paragon Plus Environment


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for the composite with no NPs. Comparing that sample to the composites with NPs, there was a minimum increase of 130% in tan δ to ca. 0.29 and a maximum increase of 257% in tan δ to ca. 0.45 for the 40 wt.% and 20 wt.% NP composites, respectively. The 1:40 composites (Figure 5d,e) experienced significant increases in the peak tan δ as well. Without NPs, the tan δ was ca. 0.14. With NP integration, the minimum tan δ increase was 65% to ca. 0.23, and the maximum increase was 147% to a value of ca. 0.35 for the 30 wt.% and 10 wt.% composites, respectively. These values are tabulated in Supporting Information Table S5 and Table S6. Despite there being no clear correlation between tan δ and NP concentration, every composite with NPs experienced a significant increase in the peak tan δ over the composite without NPs. While the peak tan δ showed significant increases for all composites with NPs, the tan δ changes before the onset of the glass transition were not as drastic. The 1:10 composites (Figure 5c) showed an initial tan δ of ca. 0.023 with no NPs but decreased slightly to ca. 0.022 with 10 wt.% NPs. With additional NPs, the tan δ increased modestly to a maximum of ca. 0.027 for the 50 wt.% NPs, which equates to a 17.0% increase. However, the 1:40 composites (Figure 5f) experienced an opposite trend in tan δ value before the glass transition temperature. For the composite with no NPs, the tan δ was ca. 0.027. This initially remained unchanged with the addition of 10 wt.% NPs. With additional NPs, the tan δ decreased to a minimum value of ca. 0.021 for the 40 wt.% NP composite. This equates to a 21% decrease in tan δ as compared to the composite with no NPs. These tan δ values can be seen in Supporting Information Table S7 and Table S8. Therefore, before the glass transition temperature, the different epoxy dilutions had experienced differing changes in the tan δ. However, at the glass transition temperature, the NP content generated profound increases in the tan δ for both epoxy dilutions.


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These dramatic changes in the tan δ at the glass transition temperature are attributed to mechanical rather than chemical contributions, which agrees with the previous sections. Just as the fibers will create an interphase region with the matrix, the NPs themselves are surrounded by an interphase layer due to polymer-chain immobilization, which can contribute to the energy absorption ability of the composite.49,

52

When elevating the temperature of the composite

through its glass transition temperature, an increase in tan δ beyond that of just the polymer will occur with the onset of slippage between the NPs and the polymer matrix when no covalent bonding is present between the NP and polymer.53, 54 Using tan δ as the curve for determining the glass transition temperature, the transition temperature was not significantly affected by the NPs at these loading concentrations thus also confirming the lack of chemical bonding between the NPs and polymer. This explains the significant increases in tan δ at the glass transition temperature for the composites with NPs.



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Figure 5. The tan δ plots for different SiC concentrations and two different epoxy dilutions: (a-c) 1:10 dilution and (d-f) 1:40 dilution. (a) Shows the full temperature range of the 1:10 dilution with magnified view (b) at the glass transition temperature and (c) before the glass transition temperature. (d) Shows the full temperature range of the 1:40 dilution with magnified view (e) at the glass transition temperature and (f) before the glass transition temperature. 18 ACS Paragon Plus Environment

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3. Conclusion Here, it was demonstrated that through a facile fiber coating process nanomaterials could be adhered to the fiber’s surface and enhance the bulk composite’s properties. SiC NP integration resulted in enhanced SHM capabilities with a maximum sensitivity increase of 53.6%, and mechanical performance enhancements seen through interlaminar strengths increased by up to 13.7%. Ideally, a multifunctional composite will have both increased mechanical strength and enhanced functional properties. The 40 wt.% SiC in 1:40 dilution composite simultaneously experienced a 47.5% increase in gauge factor and a 7.7% increase in interlaminar strength, and the 50 wt.% SiC in 1:40 dilution composite revealed a 15.8% increase in gauge factor with a 12.2 % increase in interlaminar strength. Additionally, at the glass transition temperature, all composites with SiC NPs showed substantial increases in tan δ, which ranged from a 65% to 257% increase. These results show that simple integration of low weight fractions of ceramic NPs in the fiber sizing can have profound effects on various properties of the bulk composite. Most importantly, the interlaminar strength and SHM sensitivity were enhanced simultaneously by using SiC NPs thus demonstrating a truly multifunctional composite using an easily commercializable process that is economically viable using current commercially-available nanomaterials. 4. Experimental Methods 4.1 Nanoparticle Deposition SiC NPs with a dimeter of ca. 45-65 nm were supplied by US Research Nanomaterials, Inc.. These NPs were supplier-certified as 90% beta phase and 10% amorphous. The epoxy sizing was Hydrosize® EP876 supplied by Michelman. Deionized water was added to the EP876 and 19 ACS Paragon Plus Environment


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mechanically stirred for 15 minutes at room temperature. Two different epoxy dilutions were used, 1:10 and 1:40 by weight of epoxy mixture to water. The epoxy was received as a water mixture already, so the weight of that mixture was used in calculating the amount of water to add. When adding the SiC NPs, the weight fraction was calculated from the weight of the dehydrated epoxy. SiC NP concentrations ranged from 0 wt.% to 50 wt.% in 10 wt.% increments. With the specified amount of SiC NPs weighed out, it was slowly added to the stirring diluted epoxy. As an example, to prepare a 1:10 dilution with 40 wt.% SiC NPs, 8.0 g of the EP876 epoxy was combined with 80 g of DI water and 0.31 g of SiC NPs. The mixture was left to stir for 24 hours and sonicated for approximately 30 minutes after about 1 hour of stirring at room temperature. This stirring resulted in no agglomerations of NPs that were visibly observable. This mixture was then added to the dip coating bath and was continuously mechanically stirred during fiber application process. The carbon fiber used in this research was commercially-available Hexcel HexTow ® IM7. It was supplied as an unsized, 12k tow (containing 12,000 individual carbon fibers) that was surface treated by the manufacturer to improve compatibility with epoxy. This tow was fed at a rate of approximately 40 meters per hour through the dip coating bath and dried in air at approximately 120 °C for 1 minute before being spooled on a fiber creel to await composite fabrication. The visual observation of the coating was performed using a Hitachi S4800 scanning electron microscope operated at 10 kV and 10 µA. Approximate nanoparticle-to-nanoparticle distances were calculated using SEM images and measuring 25 distances between neighboring NPs. For 1:10 epoxy dilution with 20 wt% NPs and 1:40 epoxy dilution with 40 wt% NPs fibers, the average spacing was 0.21 ± 0.10 µm and 0.33 ± 0.16 µm, respectively. 4.2 Composite Fabrication: 20 ACS Paragon Plus Environment

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The compression mold for these unidirectional composites had a fixed final volume. Based on this final composite volume, the proper volume of fiber tows was added to the mold with excess amounts of epoxy. The epoxy matrix was a 100:26.4 ratio by weight of Hexion, Inc. Epon 862 and Hexion, Inc. Epikure™ Curing Agent W. Under compression, the excess epoxy could flow from the mold, and the high viscosity nature of the epoxy allowed the proper amount of epoxy to remain in the mold during the curing process. The epoxy was cured at 121°C for 4 hours, according to manufacturer specifications. The resulting composite had a dense, nonporous structure with an average density of 1.59 ± 0.028 g/cc. Two different thicknesses of composites were fabricated with the fiber volume fraction staying constant. The thicker composites were used for the short beam shear tests and were 2.85 mm thick with a width of 12.6 mm and a length of 125 mm. Thinner composites were fabricated for the SHM and DMA tests. The width and length remained constant, but the thickness was decreased to 1mm. The reduced thickness enabled increased DMA strain amplitude of the composite beams during SHM characterization. Fiber loading was calculated to achieve a 60 vol% fiber content. This equated to using 85 tows for the thicker samples and 33 tows for the thinner samples, which was based on the fiber diameter of 5.2 µm and the two containing 12,000 fibers. 4.3 Structural Health Monitoring Characterization: A single cantilever clamp in a TA Instruments Q800 Dynamic Mechanical Analyzer (DMA) was used to strain the composite beams. The beams were sectioned and polished to a thickness, width and length of 1mm, 12.5mm and 60mm, respectively. An out-of-plane through thickness electroding configuration was established by adhering wires to the composite’s surface using silver paint. The electrodes were placed 12.5 mm apart and were centered between the two DMA clamps. A Keysight 34470A digital multimeter was used in conjunction with Keysight 21 ACS Paragon Plus Environment


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BenchVue software to record the electrical resistance data. The DMA testing used a custom displacement-controlled method that induced a specified displacement to each composite. The displacement rate was 3000 µm/min to the specified displacement level then held at that displacement for roughly 10 seconds. The DMA then returned the composite beam to zero displacement at a rate of 3000 µm/min and held for another 10 seconds before displacing again. This procedure was repeated at least 10 times at different strain levels. 4.4 Fourier transform infrared spectroscopy Fourier transform infrared spectroscopy (FTIR) measurements were performed with a PerkinElmer Frontier instrument using the attenuated total reflectance method. The spectra were measured from 500 cm-1 to 4000 cm-1 using a scan speed of 1 cm s-1 and a resolution of 4 cm-1 with the baseline being subtracted for correction. 4.5 Mechanical Characterization: Interlaminar shear strength tests were performed on a MTS Alliance RT/5 tensile frame according to ASTM D2344 with the samples sectioned and polished to the approximate dimensions of approximately 2.6 mm by 6 mm by 18mm. The dimensions for each individual sample were measured before each test to attain accurate strength results. The span length was adjusted for each sample set based on the sample thickness and in accordance to ASTM 2344. The crosshead speed was 1mm/min. These tests were performed on at least 10 samples for each composite. A TA Instruments Q800 DMA was used to characterize the viscoelastic properties of the composites as a function of temperature. The composite beams were cut to approximate 22 ACS Paragon Plus Environment

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dimensions of 1 mm, 12.5 mm and 35 mm for thickness, width, and length, respectively. Using a three-point bend configuration, the DMA was operated with an oscillating 10 Hz sinusoidal waveform with a strain amplitude of 0.01% and was ramped from room temperature to 250 °C at a rate of 2 °C/min. 4.6 Thermal Analysis: A TA Instruments Q500 thermogravimetric analyzer (TGA) was used to gather the thermal degradation characteristics of the fibers as seen in the Supporting Information Figure S1. Platinum pans were used with roughly 3-5 mg of fibers for each test. High resolution mode with a dynamic temperature ramp rate was used to capture the fiber’s degradation properties. The test ran from room temperature to 850 °C at 50 °C/min in air. As a result of the mode used for these tests, the ramp rate would decrease when large weight changes were detected in order to more precisely identify transition temperatures. Air was used instead of nitrogen in order to burn off every component in the sample except for the SiC NPs. This allowed for the determination of the epoxy and SiC content relative to the entire fiber weight. ASSOCIATED CONTENT This manuscript has been authored by UT-Battelle, LLC under Contract No. DE-AC0500OR22725 with the U.S. Department of Energy. The United States Government retains and the publisher, by accepting the article for publication, acknowledges that the United States Government retains a non-exclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this manuscript, or allow others to do so, for United States Government purposes. The Department of Energy will provide public access to these results of



federally

sponsored

research

in

accordance

with

the

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DOE

Public

Access

Plan

(http://energy.gov/downloads/doe-public-access-plan). Supporting Information TGA curves are provided in the supporting information along with schematics illustrating the testing configuration for SHM and short beam shear. All tabulated data is also provided. AUTHOR INFORMATION Corresponding Authors * Email: [email protected] * Email: [email protected] ORCID Christopher C. Bowland: 0000-0002-1229-4312 Ngoc A. Nguyen: 0000-0002-0278-406X Amit K. Naskar: 0000-0002-1094-0325 Author Contributions All the authors have made contributions to the manuscript and given their approval of the final version of the manuscript. Acknowledgements Research sponsored by the Wigner Fellowship Program as part of the Laboratory Directed Research and Development Program of Oak Ridge National Laboratory, managed by UT-


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Battelle, LLC, for the U. S. Department of Energy. N. A. N. and A. K. N. also acknowledge other programmatic support from the Laboratory Directed Research and Development Program. References (1)

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