**Advance in Mechanical and Thermal Characterization of Polymer Composites**

Editor

**SD Jacob Muthu**

Basel • Beijing • Wuhan • Barcelona • Belgrade • Novi Sad • Cluj • Manchester

*Editor* SD Jacob Muthu Faculty of Engineering & Applied Science University of Regina Regina Canada

*Editorial Office* MDPI St. Alban-Anlage 66 4052 Basel, Switzerland

This is a reprint of articles from the Special Issue published online in the open access journal *Polymers* (ISSN 2073-4360) (available at: www.mdpi.com/journal/polymers/special issues/D03UN7657W).

For citation purposes, cite each article independently as indicated on the article page online and as indicated below:

Lastname, A.A.; Lastname, B.B. Article Title. *Journal Name* **Year**, *Volume Number*, Page Range.

**ISBN 978-3-0365-8653-3 (Hbk) ISBN 978-3-0365-8652-6 (PDF) doi.org/10.3390/books978-3-0365-8652-6**

© 2023 by the authors. Articles in this book are Open Access and distributed under the Creative Commons Attribution (CC BY) license. The book as a whole is distributed by MDPI under the terms and conditions of the Creative Commons Attribution-NonCommercial-NoDerivs (CC BY-NC-ND) license.

## **Contents**

### **Branka Muˇsiˇc, Nataˇsa Knez and Janez Bernard**

## **About the Editor**

### **SD Jacob Muthu**

SD Jacob Muthu is a distinguished researcher in the field of material science and engineering and is well known for his work in polymer composites and nanocomposites. With an illustrious carrier spanning a number of years, he has contributed significantly to his research field. Throughout his academic journey, Dr. Muthu has published numerous journal papers, conference proceedings, and research reports that have enriched the scientific community.

## **Preface**

With the ever-evolving innovation in nanomaterials, polymer nanocomposites have revolutionized engineering industries and technologies across the globe and unlocked new possibilities and functionality for polymers. The combination of nanomaterials with polymers has promoted a new generation of materials that exhibit exceptional mechanical, thermal, electrical and optical properties. This Reprint is a collection of journal papers focused on the various combination of nanocomposites that analyze the intricate behaviour of nanoparticles to improve the functionalities of polymers in various engineering applications.

All the papers in this Reprint have been subjected to a rigorous peer-review process before being accepted for publication. The published articles will benefit broad audiences, including students, research engineers, and professionals seeking to understand the complexities of polymer nanocomposites. Further, this Reprint will offer a wealth of knowledge and insights into applying nanocomposites in various engineering industries.

> **SD Jacob Muthu** *Editor*

## *Article* **Dynamic Impact Properties of Carbon-Fiber-Reinforced Phenolic Composites Containing Microfillers**

**Ibraheem A. Abdulganiyu \*, Oluwasegun E. Adesola, Ikechukwuka N. A. Oguocha and Akindele G. Odeshi**

Department of Mechanical Engineering, University of Saskatchewan, Saskatoon, SK S7N 5A9, Canada<EMAIL_ADDRESS>(O.E.A<EMAIL_ADDRESS>(I.N.A.O<EMAIL_ADDRESS>(A.G.O.) **\*** Correspondence<EMAIL_ADDRESS>
**Abstract:** The addition of nano- and microfillers to carbon-fiber-reinforced polymers (CFRPs) to improve their static mechanical properties is attracting growing research interest because their introduction does not increase the weight of parts made from CFRPs. However, the current understanding of the high strain rate deformation behaviour of CFRPs containing nanofillers/microfillers is limited. The present study investigated the dynamic impact properties of carbon-fiber-reinforced phenolic composites (CFRPCs) modified with microfillers. The CFRPCs were fabricated using 2D woven carbon fibers, two phenolic resole resins (HRJ-15881 and SP-6877), and two microfillers (colloidal silica and silicon carbide (SiC)). The amount of microfillers incorporated into the CFRPCs varied from 0.0 wt.% to 2.0 wt.%. A split-Hopkinson pressure bar (SHPB), operated at momentums of 15 kg m/s and 28 kg m/s, was used to determine the impact properties of the composites. The evolution of damage in the impacted specimens was studied using optical stereomicroscope and scanning electron microscope. It was found that, at an impact momentum of 15 kg m/s, the impact properties of HRJ-15881-based CFRPCs increased with SiC addition up to 1.5 wt.%, while those of SP-6877-based composites increased only up to 0.5 wt.%. At 28 kg m/s, the impact properties of the composites increased up to 0.5 wt.% SiC addition for both SP-6877 and HRJ-15881 based composites. However, the addition of colloidal silica did not improve the dynamic impact properties of composites based on both phenolic resins at both impact momentums. The improvement in the impact properties of composites made with SiC microfiller can be attributed to improvement in crystallinity offered by the α-SiC type microfiller used in this study. No fracture was observed in specimens impacted at an impact momentum of 15 kg m/s. However, at 28 kg m/s, edge chip-off and cracks extending through the surface were observed at lower microfiller addition (≤1 wt.%), which became more pronounced at higher microfiller loading (≥1.5 wt.%).

**Keywords:** particle-reinforced composites; mechanical properties; split-Hopkinson pressure bar (SHPB); impact behaviour; crystallinity; fracture

## **1. Introduction**

Carbon-fiber-reinforced phenolic composites (CFRPCs) are used widely as structural materials in aeronautical and space industries due to their excellent properties such as low density, low thermal expansion, high thermal stability, high strength and stiffness, and excellent resistance to creep, fatigue, and corrosion [1,2]. Moreover, they can be turned into carbon-carbon composites by converting the phenolic matrix into amorphous carbon through pyrolysis [3–6].

Phenolic resins have attractive properties which make them good matrix materials for fabricating carbon-fiber-reinforced composites. Compared to other polymers, phenolic resins have lower flammability and smoke emission, superior mechanical strength, higher hardness, better heat resistance, chemical resistance, and dimensional stability [7–9]. Phenolic resin material systems have been used for ballistic protective body armor since the 1960s. Phenolic resin blended with polyvinyl butyral (PVB) resin was one of the earliest matrix

**Citation:** Abdulganiyu, I.A.; Adesola, O.E.; Oguocha, I.N.A.; Odeshi, A.G. Dynamic Impact Properties of Carbon-Fiber-Reinforced Phenolic Composites Containing Microfillers. *Polymers* **2023**, *15*, 3038. https://doi.org/10.3390/ polym15143038

Academic Editor: S. D. Jacob Muthu

Received: 23 April 2023 Revised: 30 June 2023 Accepted: 12 July 2023 Published: 13 July 2023

**Copyright:** © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

materials approved for ballistic protective body armor [10]. In the early 1960s, DeBell & Richardson Inc. (Enfield, CT, USA) developed a PVB-phenolic resin system initially for nylon helmet liners. Aircraft structures experience impact damage as a result of bird strikes, turbine blade failures, and inspiration of runaway debris into the engine, with velocities of 100–150 m/s and impact energies of over 100 J [11,12].

The structural performance of polymers and fiber-reinforced polymer composites can be enhanced through the addition of filler materials such as carbon nanotube (CNT), graphene, graphene oxide (GO), nanoclay, non-metallic oxides and carbides, and metallic oxides, carbides, borides, silicates, and silicides. The modification of matrices of CFRP composites with nano- and microfillers has been reported to lead to considerable improvement in their ablation resistance [13–19], interlaminar shear strength (ILSS) [20–27], interlaminar fracture toughness (ILFT) [22,28–30], wear and tribological properties [31–33], as well as flexural properties (strength and elastic modulus) [34,35]. Xu et al. [13] studied the ablation resistance and mechanism of CF-fabric-reinforced phenolic resin composite modified with particles of tantalum disilicide (TaSi2) and reported significant improvement in the anti-ablation property of TaSi2-modified composite in comparison with the unmodified composite. At 50 wt.% TaSi2 addition, the linear ablation rate and mass ablation rate reduced by 12% and 30%, respectively. In a different study, Wang et al. [14] found that the linear ablation rate of SiC particle-modified CF-reinforced phenolic resin composite decreased with increasing SiC content up to 5 wt.%, after which it started to increase. The effect of halloysite nanoclay addition on the mechanical and thermal properties of carbon/glass-fiber-reinforced epoxy resin composites was investigated by Nagaraja et al. [20]. They found that the interlaminar shear strength (ILSS) of nanoclay-modified composites was higher than that of the unmodified composites. At 3 wt.% halloysite nanoclay content, the ILSS of the composites improved by roughly 18% over that of unmodified composites. Lyu et al. [27] reported roughly 55% and 94% increases in ILSS and flexural strength, respectively, for CF/polyetheretherketone (PEEK) composites modified with multi-walled carbon nanotube (MWCNT) when compared to unmodified CF/PEEK composites. Kostagiannakopoulou et al. [29] studied the effect of adding graphene nanoplatelets with different geometrical characteristics on the interlaminar fracture toughness (ILFT) of CFRP composites. They observed that composites containing graphene nanoplatelets with high values of aspect ratio (AR) and specific surface area (SSA) showed higher Mode I and II ILFT than the unmodified composites.

A main challenge arising from the inclusion of micro- and nanofillers in polymer-based composites is that they agglomerate during dispersion in the matrix as a result of their large SSA and van der Waals forces [36–38]. The net effect of agglomeration is that it negatively influences the end properties of the resulting composites. Moreover, pristine carbon fiber is smooth, chemically inert, hydrophobic, and possesses low surface energy. As such, bonding between it and the polymer matrix occurs mostly through the weak van der Waals forces. Some of the methods employed in dispersing fillers in polymer matrices are vigorous mixing, shear mixing, mechanical mixing (e.g., ball-milling and 3-roll milling or calendaring), ultrasonication, magnetic stirring, surface coating of fillers, and chemical and physical surface functionalization [29,36–41]. While a perfect dispersion of the filler in different polymer matrices could be very challenging and difficult to achieve, factors such as the type of the polymer matrix (thermoplastic or thermoset), dimensions and amount of microfillers, the availability of techniques and fabrication processes are considered when selecting a proper technique for filler dispersion [29,41]. The dispersion state of fillers in matrices of polymer composites plays an important role in determining their overall performance. Its evaluation can provide an insight into the relationship between the microstructure and measured properties of the composites [42,43]. The experimental techniques commonly used for evaluating dispersion state include optical microscopy, dynamic light scattering (DLS), ultraviolet-visible-infrared (UV-Vis-IR) spectroscopy, scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction, X-ray microtomography,

and small angle scattering techniques, such small angle X-ray scattering (SAXS), small angle light scattering (SALS), and small angle neutron scattering (SANS) [42–46].

A survey of the available literature shows that, whereas several research works (Zhao et al. [47], Mohsin et al. [48], Rouf et al. [49], Lomakin et al. [50], Rouf et al. [51], Li et al. [52])have been carried out on the dynamic mechanical properties of plain CFreinforced polymer composites, there is little information on the dynamic impact behaviour of microfiller-modified CFRPCs (Chihi et al. [53]) investigated using a split-Hopkinson pressure bar (SHPB) apparatus. Zhao et al. [47] used an electromagnetic SHPB to study the influence of strain rate and loading direction on the compressive failure behaviour of 2D triaxially braided CF fabric/epoxy resin composites at dynamic strain rates of 200 s−1 and 220 s−1 . The results of the SHPB tests were compared with those of quasi-static tests performed at strain rates of 10−4 s −1 and 10−2 s −1 . Both axial and transverse compressive strengths were found to increase with strain rate, which they attributed to change in the mechanisms of damage evolution. Rouf et al. [49] studied the effect of strain rate on the compressive failure characteristics of a 1D non-crimp fabric (1D-NCF) carbon fiber/epoxy composite at strain rates of 0.003/s, 0.2/s, and 325/s. The high strain rate (325/s) test was performed using a SHPB apparatus. A strong strain rate dependency of the compressive strength of the composites was reported. On the other hand, Li et al. [52] studied the strain rate dependency of mechanical properties of warp-knitted and plain weave CF fabric/epoxy resin composites. A quasi-static test was performed at 0.5/s, while dynamic compression tests were performed with a SHPB instrument at strain rates that ranged from roughly 221/s to 1337/s. A weak strain rate dependency of the dynamic compressive strength was reported. Chihi et al. [53] investigated the dynamic response and damage characteristics of CF/epoxy resin composites containing different weight fractions of CNTs (0%, 0.5%, and 2%) using the SHPB equipment. Compression SHPB tests were conducted at impact pressures ranging from 1.4 bars to 2 bars. The results showed that the addition of CNTs significantly improved the mechanical performance of CF/epoxy composites. Within the range of CNT content and impact pressure tested, the maximum compressive strength obtained increased with the impact pressure and weight fraction of CNT.

The goal of this research is to improve the impact response of carbon-fiber-reinforced phenolic composites with microfillers (colloidal silica and micron-sized silicon carbide) at varying amounts and establish a performance threshold limit of the CFRPCs with the microfiller addition. This would help usher in a new class of high-performing CFRP composites for use in high strain rate applications without suffering catastrophic failure in service. In the long run, this research is intended to develop numerical models that could predict the CFRP material performance at high strain rate loading conditions. To develop this numerical model, material data generated at high strain rates are necessary. For this purpose, the SHPB device was used in this work to obtain high strain rates.

### **2. Materials and Methodology**

### *2.1. Materials*

Figure 1 shows the flow chart used in manufacturing CFRP composites containing microfillers. A 3K, 2 × 2 twill weave (PAN) carbon fiber (CF) fabric manufactured by Fibre Glast Developments Corporation (Brookville, OH, USA), with each fabric layer measuring 0.305 mm thick, was used as the carbon fiber material. Two resole-type phenolic resins (HRJ-15881 and SP-6877) procured from SI Group (Schenectady, NY, USA) were used as matrix materials. Their compositions and properties, as provided by the manufacturer, are shown in Table 1. Two microfillers were used to modify the matrices, namely: colloidal silica (406), which was procured from West System Inc., (Bay City, MI, USA), and silicon carbide (SiC) produced by Washington Mills, (North Grafton, MA, USA). The colloidal silica (CS) comprises sand and quartz particles with an average particle size of 0.2–0.3 µm and a density of 50 g/L. The density of the SiC was 3.19 g/cm3 and its particle size range was 0.3–2.2 µm. Polyethylene glycol (PEG) (product name: CarbowaxTM PEG-400), with a density of 1.13 g/mL, which was procured from Fisher Chemical Canada, was used to disperse the colloidal silica in the phenolic resins. PEG has long molecular chains which tend to disrupt the mutual attraction between individual microfiller particles, thereby promoting their uniform dispersion in the phenolic resins [37,38]. disperse the colloidal silica in the phenolic resins. PEG has long molecular chains which tend to disrupt the mutual aĴraction between individual microfiller particles, thereby promoting their uniform dispersion in the phenolic resins [37,38].

Polymers 2023, 15, x FOR PEER REVIEW 4 of 35

was 0.3–2.2 µm. Polyethylene glycol (PEG) (product name: CarbowaxTM PEG-400), with a density of 1.13 g/mL, which was procured from Fisher Chemical Canada, was used to

Figure 1. A flow chart showing the procedure for fabricating microfiller-modified CFRP composites. **Figure 1.** A flow chart showing the procedure for fabricating microfiller-modified CFRP composites.

Table 1. Compositions and properties of the two resole phenolic resins (HRJ-15881 and SP-6877). **Table 1.** Compositions and properties of the two resole phenolic resins (HRJ-15881 and SP-6877).

2.1.1. Dispersion of Microfillers in the Resole Phenolic Resins

2.1.1. Dispersion of Microfillers in the Resole Phenolic Resins A detailed description of the dispersion of colloidal silica (CS) and SiC microfillers in the phenolic resins can be found in the work of Abdulganiyu et al. [54]. To provide further context, CS was mixed with PEG in a ratio of 2:1 to form a shear-thickening fluid (STF). By shear thickening, we mean an increase in viscosity with increasing rate of shear strain during mechanical deformation [55]. The PEG and CS mixture was stirred using a magnetic stirrer at 1200 rpm for 3 h at room temperature. Afterward, the resulting STF was diluted with ethyl alcohol using a ratio of 3:2 and held for 2 h. The mixture was then added to the phenolic resins in the appropriate amount to obtain 0.5, 1.0, 1.5, and 2.0 wt.% of CS in the resins. To promote CS dispersal in the phenolic resins, the mixture was stirred magnetically at 1200 rpm for 1 h, followed by ultrasonication for 1 h using a Branson 1510 ultrasonicator. The ultrasonication was performed in steps of 10 min with 5 min of rest in A detailed description of the dispersion of colloidal silica (CS) and SiC microfillers in the phenolic resins can be found in the work of Abdulganiyu et al. [54]. To provide further context, CS was mixed with PEG in a ratio of 2:1 to form a shear-thickening fluid (STF). By shear thickening, we mean an increase in viscosity with increasing rate of shear strain during mechanical deformation [55]. The PEG and CS mixture was stirred using a magnetic stirrer at 1200 rpm for 3 h at room temperature. Afterward, the resulting STF was diluted with ethyl alcohol using a ratio of 3:2 and held for 2 h. The mixture was then added to the phenolic resins in the appropriate amount to obtain 0.5, 1.0, 1.5, and 2.0 wt.% of CS in the resins. To promote CS dispersal in the phenolic resins, the mixture was stirred magnetically at 1200 rpm for 1 h, followed by ultrasonication for 1 h using a Branson 1510 ultrasonicator. The ultrasonication was performed in steps of 10 min with 5 min of rest in between to prevent overheating.

between to prevent overheating. Firstly, SiC particles were weighed in the appropriate proportions to correspond to 0.5, 1.0, 1.5, and 2.0 wt.%. Then, they were mixed with 5 mL ethanol absolute to produce composites containing 0.5 wt.% SiC. The volume of ethanol mixed with SiC particles was Firstly, SiC particles were weighed in the appropriate proportions to correspond to 0.5, 1.0, 1.5, and 2.0 wt.%. Then, they were mixed with 5 mL ethanol absolute to produce composites containing 0.5 wt.% SiC. The volume of ethanol mixed with SiC particles was increased to 6, 7, and 8 mL, for composites containing 1.0, 1.5, and 2.0 wt.% SiC particles, respectively. The main reason for using ethanol absolute was to help disperse SiC particles in the phenolic resin. In addition, it was anticipated that the ethanol would evaporate when

the resole resin is heated during the curing process, leaving well-dispersed SiC particles in the phenol resin matrix of the resulting CFRP composite. The SiC particles-ethanol mixture was then ultrasonicated for a total of 30 min using the Branson 1510 ultrasonicator. The sonication was also carried out in time steps of 10 min with 5 min rest in between to prevent overheating. The ultrasonicated mixture of ethanol and SiC microfillers was then mixed with the phenolic resin and the new mixture was further sonicated for 1 h (in time steps of 10 min with 5 min rest in between to prevent overheating). Finally, after the ultrasonication, the mixture was magnetically stirred at 1200 rpm for 1 h.

### 2.1.2. Fabrication of CFRP Composites

An in-house mold measuring 170 mm × 60 mm × 10 mm (*l* × *w* × *t*) was used to manufacture CFRP composite specimens used for dynamic impact testing. The carbon fiber fabric was cut into rectangular sheets measuring 160 mm × 60 mm (*l* × *w*). Since the thickness of the internal cavity of the mold is 10 mm and thickness of the carbon fiber fabric sheet is 0.305 mm, the number of reinforcement sheets used (*M*) was calculated using Equation (1).

$$M = \frac{\text{thickness of mold internal cavity}}{\text{thickness of carbon fiber fabric}} = \frac{10 \text{ mm}}{0.305 \text{ mm}} \tag{1}$$

The internal cavity of the mold was sprayed with a SLIDE® Epoxease mold release agent (No. 40614N) produced by Infotrac (USA) and allowed to dry for 25 min. The lubricant made the removal the CFRPCs from the mold after fabrication easy. Using the hand lay-up technique, 33 sheets of the carbon fiber fabric were pressed together to fit into the 10 mm thickness of the internal cavity of the mold. Afterward, the previously prepared mixture of microfillers and phenolic resin was poured into the mold cavity and it slowly impregnated the carbon fiber fabric sheets with the help of gravity.

Curing of the resin-impregnated carbon fiber fabric was carried out in a Parr 4848 pressure reactor autoclave under an argon gas atmosphere at a pressure of 50 bar. The temperature of the autoclave was ramped up from room temperature to 120 ◦C at a heating rate of 2 ◦C/min and maintained at 120 ◦C for 1 h, after which it was cooled to room temperature. The argon gas atmosphere helped to remove gaseous by-products of polymerization reactions that occurred during curing. It also helped to reduce the porosity of the CFRPCs.

### *2.2. Characterization of Fabricated Materials*

### 2.2.1. X-ray Diffraction (XRD)

Specimens used for XRD experiment were prepared by dipping CFRP composites in liquid nitrogen and crushing them in a clean ceramic crucible to fine aggregates. XRD analysis was obtained with a Rigaku Ultima IV X-Ray Diffractometer (Rigaku Americas Corporation, The Woodlands, TX, USA) using Cu Kα source (*λ* = 1.54056 Å), tube voltage of 40 kV and tube current of 44 mA. Experimental data were collected using a Multipurpose Attachment, with para-focusing mode. A nickel Kβ filter was placed at the receiving end. X-ray diffractograms were acquired in 2θ from 5◦ to 60◦ at a scan rate of 0.5◦/min and a step size of 0.02. The software used for background correction and data smoothening was PANalytical X'Pert Highscore, while Origin 2019 (OriginLab, Northampton, MA, USA) software was used for plotting and analyzing the obtained XRD data.

### 2.2.2. Dynamic Impact Test

The high strain rate behaviour of fabricated CFRP composites was investigated using a split-Hopkinson pressure bar (SHPB) apparatus, the schematic diagram of which is presented in Figure 2. As can be seen, it comprises three bars, namely: striker, incident, and transmitter bars. The transmitter and incident bars were machined from aluminum alloy 7075-T651 and have a diameter of 38 mm. Figure 3 shows a photographic image of typical impact test specimens, each measuring 12 mm × 12 mm × 10 mm. The surfaces

of test specimens in contact with the incident and the transmitted bars were lubricated with Vaseline to reduce friction between the contacting end surfaces. With the specimen placed along the thickness between the incident and transmitted bars, the striker bar was fired with the aid of the gas gun toward the incident bar at pressures of 50 kPa and 80 kPa which corresponded to impact momentums of 15 kg m/s and 28 kg m/s, respectively. Five specimens of each CFRP composite and neat cured resin were tested at each impact momentum. Polymers 2023, 15, x FOR PEER REVIEW 7 of 35 The two firing pressures used in this study, 50 kPa and 80 kPa, generated average true strain rates of 880 s−1 and 2150 s−1 in the test specimens, respectively.

Figure 2. A schematic diagram of the split-Hopkinson pressure bar (SHPB) system. **Figure 2.** A schematic diagram of the split-Hopkinson pressure bar (SHPB) system.

Figure 3. Photographic image showing typical dynamic compression test specimens. **Figure 3.** Photographic image showing typical dynamic compression test specimens.

2.2.3. Microstructure Examination Specimens with dimensions 10 mm × 10 mm × 10 mm, cut from the fabricated rectangular plates of CFRP composites using a Buehler abrasive cuĴer (Model 95, C1800), were used for microstructure analysis. They were ground initially using 320, 500, 800, and 1200 SiC grit emery papers, followed by fine grinding using 2000 and 4000 SiC grit emery pa-During SHPB tests, an incident stress pulse is produced when the striker bar impacts the incident bar which is propagated along the incident bar toward the specimen–incident bar interface. A portion of this stress pulse is back-reflected through the incident bar, while the remainder is propagated to the transmitter bar through the specimen. Since a constant strain rate helps to eliminate the radial inertia effects on test specimens during SHPB

pers and finally polished using a 5 µm MD-Dac cloth with 5 µm MD-Dac suspension and

stereomicroscope, Melville, NY, USA, interfaced with a PAX-it digital camera) and scanning electron microscopy (SEM, JEOL-JSM-6010LV, Peabody, MA, USA). To reduce charging and improve the image quality, SEM specimens were gold-coated using an Edward S150B spuĴer coater (BOC Edwards, Crawley, UK). The SEM was operated at an accelerating voltage of 20 kV and images were acquired in the secondary electron (SE) imaging

To properly index the X-ray diffraction (XRD) paĴerns obtained experimentally for microfillers and carbon fiber (CF), they were compared to standard diffraction peaks. Indexing of experimental diffraction peaks of the CF was carried out using the Mercury software (version 4.3.0) and graphite crystallographic information (.cif) file (deposition number: 918549). The graphite cif data were sourced from the crystal structure database of the Cambridge crystallographic data center (CCDC). The simulation of theoretical diffraction paĴerns of graphite was performed using full width at half-maximum (FWHM) values of 1.0, 2.0, 3.0, 4.0, and 5.0. The diffraction paĴerns generated using these FWHM values are shown in Figure 4. Clearly, the characteristic intense peak at roughly 2θ = 26.2° (002) is followed by minor peaks at 2θ~44.0° (101) and 2θ~54.0° (004). The existence of additional peaks at 2θ~42.3° (100) and 2θ~50.5° (102) depends on the FWHM value used.

mode.

3.1. XRD

3. Results and Discussion

tests [56], a pulse shaper with a thickness of 12 mm but the same diameter as the incident and the transmitter bars was attached to the end of the incident bar. As such, on firing at a given impact momentum, the striker came into contact with the pulse shaper to help achieve a constant strain rate. Strain gauges were used to capture the generated elastic strain waves (incident, reflected, and transmitted) traveling through the incident and transmitter bars. Amplification and conditioning of the elastic strain waves were carried out using the connected strain conditioner-amplifier system. The output from the conditioner-amplifier system was linked to the mixed-signal digital oscilloscope which, in turn, was connected to the computer. The oscilloscope output was controlled and graphically displayed by in-house developed LabView software program. Based on stress equilibrium and obtained experimental data, the engineering or nominal stress (σe), engineering strain (εe), and initial engineering strain rate . εe were calculated using Equations (2)–(4), respectively [57]:

$$
\sigma\_{\mathbf{e}} = \frac{A\_b}{A\_s} E\_b \varepsilon\_T \tag{2}
$$

$$\varepsilon\_{\mathbf{e}} = -2\frac{\mathbf{C}\_b}{L\_{\mathbf{s}}} \int\_0^\mathbf{t} \varepsilon\_R d\mathbf{t} \tag{3}$$

$$
\dot{\varepsilon}\_{\text{e}} = -2 \frac{\mathcal{C}\_b}{L\_{\text{s}}} \varepsilon\_R \tag{4}
$$

where *AS* and *Ab* stand for cross-sectional areas of the specimen and bars, respectively; ε*T* and ε*R* stand for transmitted and reflected strain waves, respectively; *Cb* is the elastic wave velocity in the bars; *EB* stands for the Young's modulus of the bar materials; while *Ls* and t represent the original length of the specimen and deformation time, respectively. True stress (σt) vs. true strain (εt) curves were generated based on Equations (5) and (6), while the true strain rate . εt was calculated using Equation (7) [58].

$$
\sigma\_l = \sigma\_\mathbf{e} (1 - \varepsilon\_\mathbf{e}) \tag{5}
$$

$$
\varepsilon\_{\rm t} = -\ln(1 - \varepsilon\_{\rm e}) \tag{6}
$$

$$
\dot{\varepsilon}\_{\text{f}} = \frac{\dot{\varepsilon}\_{\text{e}}}{(1 - \varepsilon\_{\text{e}})} \tag{7}
$$

The two firing pressures used in this study, 50 kPa and 80 kPa, generated average true strain rates of 880 s−1 and 2150 s−1 in the test specimens, respectively.

.

### 2.2.3. Microstructure Examination

Specimens with dimensions 10 mm × 10 mm × 10 mm, cut from the fabricated rectangular plates of CFRP composites using a Buehler abrasive cutter (Model 95, C1800), were used for microstructure analysis. They were ground initially using 320, 500, 800, and 1200 SiC grit emery papers, followed by fine grinding using 2000 and 4000 SiC grit emery papers and finally polished using a 5 µm MD-Dac cloth with 5 µm MD-Dac suspension and a 1 µm MD-Nap cloth with 1 µm MD-Nap suspension. Microstructures of the polished and fractured specimens were analyzed using optical microscopy (OM, Nikon Optihot stereomicroscope, Melville, NY, USA, interfaced with a PAX-it digital camera) and scanning electron microscopy (SEM, JEOL-JSM-6010LV, Peabody, MA, USA). To reduce charging and improve the image quality, SEM specimens were gold-coated using an Edward S150B sputter coater (BOC Edwards, Crawley, UK). The SEM was operated at an accelerating voltage of 20 kV and images were acquired in the secondary electron (SE) imaging mode.

### **3. Results and Discussion**

### *3.1. XRD*

To properly index the X-ray diffraction (XRD) patterns obtained experimentally for microfillers and carbon fiber (CF), they were compared to standard diffraction peaks. Indexing of experimental diffraction peaks of the CF was carried out using the Mercury software (version 4.3.0) and graphite crystallographic information (.cif) file (deposition number: 918549). The graphite cif data were sourced from the crystal structure database of the Cambridge crystallographic data center (CCDC). The simulation of theoretical diffraction patterns of graphite was performed using full width at half-maximum (FWHM) values of 1.0, 2.0, 3.0, 4.0, and 5.0. The diffraction patterns generated using these FWHM values are shown in Figure 4. Clearly, the characteristic intense peak at roughly 2θ = 26.2◦ (002) is followed by minor peaks at 2θ~44.0◦ (101) and 2θ~54.0◦ (004). The existence of additional peaks at 2θ~42.3◦ (100) and 2θ~50.5◦ (102) depends on the FWHM value used. Polymers 2023, 15, x FOR PEER REVIEW 9 of 35

Figure 4. Diffraction paĴerns of graphite computed using different values of FWHM. **Figure 4.** Diffraction patterns of graphite computed using different values of FWHM.

The calculated diffraction paĴerns were compared with the diffraction paĴern obtained experimentally for the carbon fiber used in the present study (designated as Exp. CF). It was found that the XRD paĴern calculated using an FWHM value of 4.0 matched the XRD paĴern obtained experimentally for CF (Exp. CF). The two XRD paĴerns are coploĴed in Figure 5 for comparison. For Exp. CF, it can be seen that the characteristic peaks which appear at 2θ~25.7° (002), 2θ~44.1° (101), and 2θ~54° (004) are consistent with those obtained for the XRD paĴern calculated using an FWHM value of 4.0. Nevertheless, a close look at the two diffraction paĴerns shows that the calculated (002) peak differs by 0.5° from that of Ex. CF, while the difference is 0.1° for the (101) peak. The calculated diffraction patterns were compared with the diffraction pattern obtained experimentally for the carbon fiber used in the present study (designated as Exp. CF). It was found that the XRD pattern calculated using an FWHM value of 4.0 matched the XRD pattern obtained experimentally for CF (Exp. CF). The two XRD patterns are co-plotted in Figure 5 for comparison. For Exp. CF, it can be seen that the characteristic peaks which appear at 2θ~25.7◦ (002), 2θ~44.1◦ (101), and 2θ~54◦ (004) are consistent with those obtained for the XRD pattern calculated using an FWHM value of 4.0. Nevertheless, a close look at the two diffraction patterns shows that the calculated (002) peak differs by 0.5◦ from that of Ex. CF, while the difference is 0.1◦ for the (101) peak.

Figure 5. Diffraction paĴern of graphite computed using an FWHM value of 4.0 compared with the experimental diffraction paĴern obtained for carbon fiber (Exp. CF). **Figure 5.** Diffraction pattern of graphite computed using an FWHM value of 4.0 compared with the experimental diffraction pattern obtained for carbon fiber (Exp. CF).

SiC exists in different polymorphs with α-SiC (6H hexagonal) and β-SiC (3C-cubic) being the most stable structures [59,60]. While α-SiC possesses a hexagonal crystal structure and is made at temperature higher than 1700 °C, β-SiC has a cubic crystal structure and is prepared at temperatures lower than 1700 °C. Moissanite occurs naturally as SiC with the 6H hexagonal polymorph [61]. As such, it was used as a standard material with which to compare the SiC microfiller used in this study. The cif file of moissanite used in the present study was retrieved from the American mineralogy crystal structure database (AMCSD). The Mercury software was used to index the diffraction paĴern of moissanite generated with an FWHM value of 0.6. The reason why 0.6 was used was because, as shown in Figure 6, the resulting diffraction paĴern looked closest to that of the SiC microfiller (designated as Exp. SiC) used in modifying the matrix of CFRP composites developed in the present study. The characteristic features of the diffraction paĴern of moissanite (α–SiC) (black colour) are diffraction peaks at 2θ~34.2° (101), 35.7° (102), 38.3° (103), 41.4° (104), 45.3° (105), 54.7° (107), and 60.9° (108). Several of these peaks appeared in the diffraction paĴern of Exp. SiC (red colour), indicating that the SiC microfiller used is of SiC exists in different polymorphs with α-SiC (6H hexagonal) and β-SiC (3C-cubic) being the most stable structures [59,60]. While α-SiC possesses a hexagonal crystal structure and is made at temperature higher than 1700 ◦C, β-SiC has a cubic crystal structure and is prepared at temperatures lower than 1700 ◦C. Moissanite occurs naturally as SiC with the 6H hexagonal polymorph [61]. As such, it was used as a standard material with which to compare the SiC microfiller used in this study. The cif file of moissanite used in the present study was retrieved from the American mineralogy crystal structure database (AMCSD). The Mercury software was used to index the diffraction pattern of moissanite generated with an FWHM value of 0.6. The reason why 0.6 was used was because, as shown in Figure 6, the resulting diffraction pattern looked closest to that of the SiC microfiller (designated as Exp. SiC) used in modifying the matrix of CFRP composites developed in the present study. The characteristic features of the diffraction pattern of moissanite (α–SiC) (black colour) are diffraction peaks at 2θ~34.2◦ (101), 35.7◦ (102), 38.3◦ (103), 41.4◦ (104), 45.3◦ (105), 54.7◦ (107), and 60.9◦ (108). Several of these peaks appeared in the diffraction pattern of Exp. SiC (red colour), indicating that the SiC microfiller used is of the α–SiC type.

the α–SiC type. For colloidal silica (CS), its cif file could not be found to generate the calculated XRD pattern with which to compare the diffraction pattern of the colloidal silica used in this study (Exp. CS). However, for the Exp. CS, the indexing of the peaks was obtained and compared with values from the literature. Munasir et al. [62] and Jeon et al. [63] obtained (101) peaks for colloidal silica at 2θ~21.5◦ using the International Center for Diffraction Data (ICDD) PDF-2 database (PDF-2 No. 01-087-2096) and Inorganic Crystal Structure Database (ICSD) as standards. In the present study, a prominent peak at 2θ~21.76◦ was obtained for the colloidal silica corresponding to the (101) peak from the literature. We can infer that these values agree with one another, even though there is a 0.2◦ difference between the values reported in the literature and those obtained in this study. The corresponding XRD pattern of the colloidal silica microfiller is shown in corresponding XRD diffractograms, in relation to when colloidal silica microfiller is used.

Figure 6. Calculated XRD paĴern of α-SiC compared with that of the SiC microfiller (Exp. SiC). **Figure 6.** Calculated XRD pattern of α-SiC compared with that of the SiC microfiller (Exp. SiC).

For colloidal silica (CS), its cif file could not be found to generate the calculated XRD paĴern with which to compare the diffraction paĴern of the colloidal silica used in this study (Exp. CS). However, for the Exp. CS, the indexing of the peaks was obtained and compared with values from the literature. Munasir et al. [62] and Jeon et al. [63] obtained (101) peaks for colloidal silica at 2θ~21.5° using the International Center for Diffraction Data (ICDD) PDF-2 database (PDF-2 No. 01-087-2096) and Inorganic Crystal Structure Database (ICSD) as standards. In the present study, a prominent peak at 2θ~21.76° was ob-A comparison of XRD patterns obtained for CFRPCs made from phenolic resins HRJ-15881 and SP-6877 with those obtained for CF and monolithic resin is presented in Figure 7. Two conspicuous peaks at 2θ~18.8◦ and 25.7◦ for HRJ-15881 resin, and 2θ~18.6◦ and 25.7◦ for SP-6877 resin resulting from the addition of CF to the resins, can be seen. Obviously, this is due to the fact that the identities of the constituent materials were unaffected by the manufacturing process used in this study. Furthermore, it is seen that the intensity of the characteristic (002) peak of SP-6877 resin decreased, probably due to its low viscosity.

tained for the colloidal silica corresponding to the (101) peak from the literature. We can infer that these values agree with one another, even though there is a 0.2° difference between the values reported in the literature and those obtained in this study. The corresponding XRD paĴern of the colloidal silica microfiller is shown in corresponding XRD diffractograms, in relation to when colloidal silica microfiller is used. A comparison of XRD paĴerns obtained for CFRPCs made from phenolic resins HRJ-15881 and SP-6877 with those obtained for CF and monolithic resin is presented in Figure 7. Two conspicuous peaks at 2θ~18.8° and 25.7° for HRJ-15881 resin, and 2θ~18.6° and 25.7° for SP-6877 resin resulting from the addition of CF to the resins, can be seen. Obviously, this is due to the fact that the identities of the constituent materials were unaffected by the manufacturing process used in this study. Furthermore, it is seen that the intensity of the characteristic (002) peak of SP-6877 resin decreased, probably due to its low viscosity. Figures 8 and 9 show, respectively, XRD patterns obtained for CFRPCs modified with SiC and CS microfillers for HRJ-15881 and SP-6877 resins. The addition of SiC filler to CFRP composites made with the two phenolic resins caused the intensity of the (002) peak to decrease. Sekhar and Varghese [64] investigated the thermal, mechanical, and rheological properties of phenolic resin modified with intercalated graphite bisulfate and found that the characteristic diffraction peak of the graphite moved to the left and its intensity reduced in comparison to the characteristic peak of natural graphite. Based on these results, it was concluded that the intercalated graphite truly intercalates and contains bisulfate. Also, Ki Park et al. [65] investigated the structural and electrochemical evolution of structured V2O5 microspheres during Li+ intercalation and found progressive changes in the characteristic diffraction peak of V2O5 as phase transformation occurred during Li+ intercalation. The characteristic (001) peak of V2O5 shifted left to a lower 2θ angle. The shift was attributed to α-V2O5 transformation to LixV2O5 at the onset of Li+ intercalation. In the present study, the decrease in the intensity of the (002) peak of SiC-modified phenolic resin could be due to the penetration of the phenolic matrix by SiC microfillers. Furthermore, the characteristic (102) peak of α-SiC occurred at 35.7◦ for CFRPCs made with SiC-modified HRJ-15881 and SP-6877 phenolic resins.

Figure 7. XRD paĴerns obtained for carbon fiber, monolithic resins, and their composites: (a) HRJ-15881 resin and its composite and (b) SP-6877 resin and its composite. **Figure 7.** XRD patterns obtained for carbon fiber, monolithic resins, and their composites: (**a**) HRJ-15881 resin and its composite and (**b**) SP-6877 resin and its composite. Polymers 2023, 15, x FOR PEER REVIEW 13 of 35

Figure 8. XRD diffraction paĴerns of CFRP composites modified with 0.5 wt.% SiC microfiller: (a) based on HRJ-15881 resin and (b) based on SP-6877 resin. **Figure 8.** XRD diffraction patterns of CFRP composites modified with 0.5 wt.% SiC microfiller: (**a**) based on HRJ-15881 resin and (**b**) based on SP-6877 resin.

matrices of the CF composites manufactured using both resins [64,65].

Similar to the results presented in Figure 8, it is seen in Figure 9 that CS addition to the CFRP composites caused the intensity of the characteristic (002) peaks of HRJ-15881

Figure 9. XRD paĴerns of CFRP composites modified with 0.5 wt.% CS: (a) HRJ-15881 resin and (b) SP-6877 resin. **Figure 9.** XRD patterns of CFRP composites modified with 0.5 wt.% CS: (**a**) HRJ-15881 resin and (**b**) SP-6877 resin.

Figures 10 and 11, respectively, show changes in the intensity and angular position of XRD paĴerns of neat HRJ-15881 and SP-6877 resins due to the addition of SiC and CS microfillers. Figure 10 shows that SiC microfiller addition (1.0 wt.%) resulted in the existence of a (002) diffraction peak at 2θ~18.4° for both resins. The peaks at 2θ~34.2° (101), Similar to the results presented in Figure 8, it is seen in Figure 9 that CS addition to the CFRP composites caused the intensity of the characteristic (002) peaks of HRJ-15881 and SP-6877 resins to reduce, which is an indication of the presence of CS in the respective matrices of the CF composites manufactured using both resins [64,65].

35.7° (102), and 38.3° (103) are characteristic features of α–SiC as shown previously in Figure 8, which confirms that SiC microfillers are present and intercalated in the phenolics [64,65]. Figure 11 shows the effect of adding 1.0 wt.% CS to both resins. The prominent peak at 2θ~18.2° (green colour) shifted slightly to the left when compared with the characteristic peaks of individually cured resins which appeared at 2θ~18.8° and 2θ~18.6° for HRJ-15881 and SP-6877 resins, respectively. The peak shift and decrease in intensity due to CS addition may be due to its intercalation in the phenolic resins. Therefore, on the basis of the XRD paĴerns shown in Figure 11, it could be concluded that CS fillers are present and intercalated in both phenolic matrices [64,65]. Figures 10 and 11, respectively, show changes in the intensity and angular position of XRD patterns of neat HRJ-15881 and SP-6877 resins due to the addition of SiC and CS microfillers. Figure 10 shows that SiC microfiller addition (1.0 wt.%) resulted in the existence of a (002) diffraction peak at 2θ~18.4◦ for both resins. The peaks at 2θ~34.2◦ (101), 35.7◦ (102), and 38.3◦ (103) are characteristic features of α–SiC as shown previously in Figure 8, which confirms that SiC microfillers are present and intercalated in the phenolics [64,65]. Figure 11 shows the effect of adding 1.0 wt.% CS to both resins. The prominent peak at 2θ~18.2◦ (green colour) shifted slightly to the left when compared with the characteristic peaks of individually cured resins which appeared at 2θ~18.8◦ and 2θ~18.6◦ for HRJ-15881 and SP-6877 resins, respectively. The peak shift and decrease in intensity due to CS addition may be due to its intercalation in the phenolic resins. Therefore, on the basis of the XRD patterns shown in Figure 11, it could be concluded that CS fillers are present and intercalated in both phenolic matrices [64,65].

### *3.2. Dynamic Impact Properties*

### 3.2.1. Impact Properties Obtained at a Momentum of 15 kg m/s

Figures 12–16 present results obtained from SHPB tests carried out at an impact momentum (IM) of 15 kg m/s. Figure 12 shows the true stress–true strain curves obtained for unmodified CFRP composites. Figures 13 and 14 show, respectively, the true stresstrue strain and maximum true stress plots obtained for SiC-modified CFRP composites. Similarly, Figures 15 and 16 present the true stress–true strain and maximum true stress plots obtained for CS-modified CFRP composites, respectively.

Figure 10. XRD paĴerns of phenolic resins modified with 1 wt.% SiC microfiller: (a) HRJ-15881 and (b) SP-6877. **Figure 10.** XRD patterns of phenolic resins modified with 1 wt.% SiC microfiller: (**a**) HRJ-15881 and (**b**) SP-6877.

As shown in Figure 12, impact strength and toughness of the unmodified CFRP composites fabricated with HRJ-15881 phenolic resin are better than those of unmodified composites made with SP-6877 resin. A close examination of Figure 13 shows that up to 1.5 wt.% SiC addition, the CFRP composites based on HRJ-15881 matrix showed better impact strength than the unmodified CFRP composites. However, the impact strength of composites fabricated with SP-6877 resin improved over those of unmodified composites with 0.5 wt.% SiC addition, beyond which it decreased with increasing SiC content.

A plot of maximum true stresses obtained for CFRP composites modified with SiC microfiller is shown in Figure 14. The observed trend is similar to that of the true stress–true strain curves of Figure 13. It can be seen that the impact strength of CFRP composites based on HRJ-15881 matrix is higher than the impact strengths of those made with unmodified and SiC filler-modified SP-6877 phenolic matrix.

Figure 15 shows that the impact strength of CFRP composites modified with CS filler generally deteriorated with increasing microfiller content except for composites based on HRJ-15881 matrix with 1.5 wt.% and 2.0 wt.% CS. The cause of this discrepancy in impact strength within this range of CS content is not clear. The observed trend of reduction in impact strength with increasing CS content could be attributed to particle agglomeration. Abdulganiyu et al. [54] observed a deterioration in flexural properties with colloidal silica addition to CFRP composites owing to the agglomeration of colloidal silica particles. Nevertheless, a close look at Figure 16 indicates that the impact strength of CFRP composites based on HRJ-15881 matrix is higher than those based on SP-6877 matrix at all CS filler contents.

Figure 11. XRD diffraction paĴerns obtained for phenolic resins modified with 1 wt.% CS: (a) HRJ-15881 and (b) SP-6877. **Figure 11.** XRD diffraction patterns obtained for phenolic resins modified with 1 wt.% CS: (**a**) HRJ-15881 and (**b**) SP-6877. Polymers 2023, 15, x FOR PEER REVIEW 17 of 35

**Figure 12.** Figure 12. True stress–strain curves of unmodified CFRP composites (IM = 15 kg m/s). True stress–strain curves of unmodified CFRP composites (IM = 15 kg m/s).

Figure 13. Typical true stress–true strain curves of SiC filler-modified CFRP composites (IM = 15 kg m/s): (a) HRJ-15881 and (b) SP-6877. **Figure 13.** Typical true stress–true strain curves of SiC filler-modified CFRP composites (IM = 15 kg m/s): (**a**) HRJ-15881 and (**b**) SP-6877.

Figure 14. Comparison of maximum true stresses of SiC filler-modified CFRP composites made with the two phenolic matrices (IM = 15 kg m/s). Error bars are based on standard deviation. **Figure 14.** Comparison of maximum true stresses of SiC filler-modified CFRP composites made with the two phenolic matrices (IM = 15 kg m/s). Error bars are based on standard deviation.

3.2.2. Impact Properties Obtained at a Momentum of 28 kg m/s

Figure 15 shows that the impact strength of CFRP composites modified with CS filler generally deteriorated with increasing microfiller content except for composites based on HRJ-15881 matrix with 1.5 wt.% and 2.0 wt.% CS. The cause of this discrepancy in impact strength within this range of CS content is not clear. The observed trend of reduction in impact strength with increasing CS content could be aĴributed to particle agglomeration. Abdulganiyu et al. [54] observed a deterioration in flexural properties with colloidal silica Figures 17–21 show the results of SHPB tests performed on specimens of CFRP composites at a momentum of 28 kg m/s. Figure 17 shows the true stress vs. true strain curves obtained for unmodified CFRP composites, while Figures 18 and 19 show, respectively, the true stress–true strain curves and maximum true stresses obtained for SiC filler-modified composites. Graphs of true stress vs. true strain and maximum true stresses obtained for CS filler-modified composites are shown, respectively, in Figures 20 and 21.

addition to CFRP composites owing to the agglomeration of colloidal silica particles. Nevertheless, a close look at Figure 16 indicates that the impact strength of CFRP composites based on HRJ-15881 matrix is higher than those based on SP-6877 matrix at all CS filler

contents.

Figure 15. Typical true stress–true strain curves of CS filler-modified CFRP composites (IM = 15 kg m/s): (a) HRJ-15881 and (b) SP-6877. **Figure 15.** Typical true stress–true strain curves of CS filler-modified CFRP composites (IM = 15 kg m/s): (**a**) HRJ-15881 and (**b**) SP-6877.

Figure 16. Comparison of maximum true stresses of CS filler-modified CFRP composites made with the two phenolic matrices (IM = 15 kg m/s). Error bars are based on standard deviation. **Figure 16.** Comparison of maximum true stresses of CS filler-modified CFRP composites made with the two phenolic matrices (IM = 15 kg m/s). Error bars are based on standard deviation.

3.2.2. Impact Properties Obtained at a Momentum of 28 kg m/s Figures 17–21 show the results of SHPB tests performed on specimens of CFRP composites at a momentum of 28 kg m/s. Figure 17 shows the true stress vs. true strain curves obtained for unmodified CFRP composites, while Figures 18 and 19 show, respectively, the true stress–true strain curves and maximum true stresses obtained for SiC filler-modified composites. Graphs of true stress vs. true strain and maximum true stresses obtained for CS filler-modified composites are shown, respectively, in Figures 20 and 21. Figure 17 shows that the impact strength of unmodified CFRP composites based on HRJ-15881 matrix is higher than that based on SP-6877 matrix, which is consistent with the results of unmodified composites tested at 15 kg m/s (see Figure 12). Figures 18 and 19 show that, within the range of filler content tested in this study, the impact strength of SiC filler-modified CFRP composites based on HRJ-15881 matrix is higher than that of those based on SP-6877 resin. A close examination of the true stress–true strain curves in Figure 18 shows that modification of SP-6877 resin with SiC filler improved the impact strength of the resulting composites above that of the unmodified composite at all SiC contents. In contrast, impact strength improvement is seen in HRJ-15881-based composites at only 0.5 wt.% SiC filler addition. Further addition of SiC filler beyond this level resulted in a decrease in impact strength of the resulting composites.

Figure 17. Typical true stress–true strain curves of unmodified CFRP composites (IM = 28 kg m/s). **Figure 17.** Typical true stress–true strain curves of unmodified CFRP composites (IM = 28 kg m/s).

Figure 18. Typical true stress–true strain curves of SiC filler-modified CFRP composites (IM = 28 kg m/s): (a) HRJ-15881 and (b) SP-6877. **Figure 18.** Typical true stress–true strain curves of SiC filler-modified CFRP composites (IM = 28 kg m/s): (**a**) HRJ-15881 and (**b**) SP-6877.

Figure 19. Comparison of maximum true stresses of SiC filler-modified CFRP composites based on the two resins (IM = 28 kg m/s). Error bars are based on standard deviation. **Figure 19.** Comparison of maximum true stresses of SiC filler-modified CFRP composites based on the two resins (IM = 28 kg m/s). Error bars are based on standard deviation.

strength of CFRP composites based on HRJ-15881 resin is higher those based on SP-6877 matrix at all CS contents. Furthermore, the impact strength of composites based on SP-6877 matrix generally decreased with increasing CS content. In contrast, the decrease in the impact strength of composites based on HRJ-15881 matrix did not follow a consistent trend. For all intents and purposes, the trend of the impact strength results obtained for

As reported previously for composites tested at 15 kg m/s, Figures 20 and 21 show

composites tested at 15 kg m/s is similar to those tested at 28 kg m/s.

Figure 20. Typical true stress–true strain curves of CFRP composites modified with CS filler (IM = 28 kg m/s): (a) HRJ-15881 and (b) SP-6877. **Figure 20.** Typical true stress–true strain curves of CFRP composites modified with CS filler (IM = 28 kg m/s): (**a**) HRJ-15881 and (**b**) SP-6877.

Figure 21. Comparison of maximum true stresses of CS filler-modified CFRP composites based on the two resins (IM = 28 kg m/s). Error bars are based on standard deviation. **Figure 21.** Comparison of maximum true stresses of CS filler-modified CFRP composites based on the two resins (IM = 28 kg m/s). Error bars are based on standard deviation.

It is evident from the preceding results of SHPB tests, conducted at 15 kg m/s and 28 kg m/s, that the modification of HRJ-15881 and SP-6877 phenolic matrices with SiC filler enhanced the impact resistance of their respective CFRP composites. In contrast, modifying the resins with CS reduced the impact strength of the resulting composites. Broadly, the trends of impact strengths obtained in the present study are consistent with the previous work by Abulganiyu et al. [54] in which the flexural properties of CFRP composites improved with SiC filler addition, but deteriorated with CS addition due to the agglomeration of CS particles. As reported previously for composites tested at 15 kg m/s, Figures 20 and 21 show that modification of both resins with CS filler degraded the impact resistance of the resulting CFRP composites. Nevertheless, within the range of CS content studied, the impact strength of CFRP composites based on HRJ-15881 resin is higher those based on SP-6877 matrix at all CS contents. Furthermore, the impact strength of composites based on SP-6877 matrix generally decreased with increasing CS content. In contrast, the decrease in the impact strength of composites based on HRJ-15881 matrix did not follow a consistent trend. For all intents and purposes, the trend of the impact strength results obtained for composites tested at 15 kg m/s is similar to those tested at 28 kg m/s.

3.3. Microstructural Examination Table 2 provides an overview of a visual examination of the failure conditions of specimens of CFRP composites subjected to SHPB testing at the two momentums. At 15 kg m/s, none of the specimens, irrespective of whether they were modified with SiC and CS microfillers or unmodified, fractured. Nevertheless, at 28 kg m/s, specimens of unmodified CFRP composite fractured. Since the fracture features observed for composites based It is evident from the preceding results of SHPB tests, conducted at 15 kg m/s and 28 kg m/s, that the modification of HRJ-15881 and SP-6877 phenolic matrices with SiC filler enhanced the impact resistance of their respective CFRP composites. In contrast, modifying the resins with CS reduced the impact strength of the resulting composites. Broadly, the trends of impact strengths obtained in the present study are consistent with the previous work by Abulganiyu et al. [54] in which the flexural properties of CFRP

on the two resins were similar, only CFRP composites based on HRJ-15881 matrix would

wt.% SiC will be presented. Specimens modified with 1.5 wt.% and 2.0 wt.% SiC filler suffered no fracture due to impact loading at 28 kg m/s. For CS-modified CFRP composites, no fracture was observed below 2.0 wt.% CS content. At 2.0 wt.% CS, specimens of composites based on both resins fractured. Again, failure features were similar for composites based on both resins and, as such, subsequent discussion will focus on HRJ-15881-

Table 2. Visual inspection results obtained for CFRP composite specimens tested at two impact mo-

Momentum Filler Content Did Fracture Occur?

For SiC filler-modified CFRP composites tested at 28 kg m/s, only specimens modi-

be used as representatives to avoid duplication.

based composites.

mentums.

composites improved with SiC filler addition, but deteriorated with CS addition due to the agglomeration of CS particles.

### *3.3. Microstructural Examination*

Table 2 provides an overview of a visual examination of the failure conditions of specimens of CFRP composites subjected to SHPB testing at the two momentums. At 15 kg m/s, none of the specimens, irrespective of whether they were modified with SiC and CS microfillers or unmodified, fractured. Nevertheless, at 28 kg m/s, specimens of unmodified CFRP composite fractured. Since the fracture features observed for composites based on the two resins were similar, only CFRP composites based on HRJ-15881 matrix would be used as representatives to avoid duplication.

**Table 2.** Visual inspection results obtained for CFRP composite specimens tested at two impact momentums.

For SiC filler-modified CFRP composites tested at 28 kg m/s, only specimens modified with 0.5 and 1.0 wt.% SiC fractured. Again, as the failure features found in these two CFRP composites were similar, only the detailed microstructures of composites with 0.5 wt.% SiC will be presented. Specimens modified with 1.5 wt.% and 2.0 wt.% SiC filler suffered no fracture due to impact loading at 28 kg m/s. For CS-modified CFRP composites, no fracture was observed below 2.0 wt.% CS content. At 2.0 wt.% CS, specimens of composites based on both resins fractured. Again, failure features were similar for composites based on both resins and, as such, subsequent discussion will focus on HRJ-15881-based composites.

Figures 22–24 show typical optical images of failed specimens of CFRP composites tested at 28 kg m/s. Figure 22 shows typical surfaces of failed specimens of unmodified CFRP composite, Figure 23 shows the same in the context of a CFRP composite modified with 0.5 wt.% SiC, while Figure 24 shows the same regarding a CFRP composite modified with 2.0 wt.% CS. The failure was typified by splitting the specimens into two, with the main failure mode being delamination. In addition, the failed specimens featured undulating, rough surfaces, which indicates that the matrix and fiber bundles ruptured during testing. This type of surface characteristic was more pronounced in SiC filler-modified CFRP composites than in those modified with CS filler. The surface morphology of composites modified with CS filler and those containing no fillers was similar.

posites modified with CS filler and those containing no fillers was similar.

15 kg m/s None

15 kg m/s

28 kg m/s

15 kg m/s

28 kg m/s

28 kg m/s Yes Yes

0.5 wt. % SiC No No 1.0 wt. % SiC No No 1.5 wt. % SiC No No 2.0 wt. % SiC No No

0.5 wt. % SiC Yes Yes 1.0 wt. % SiC Yes Yes 1.5 wt. % SiC No No 2.0 wt. % SiC No No

0.5 wt. % CS No No 1.0 wt. % CS No No 1.5 wt. % CS No No 2.0 wt. % CS No No

0.5 wt. % CS No No 1.0 wt. % CS No No 1.5 wt. % CS No No 2.0 wt. % CS Yes Yes

Figures 22–24 show typical optical images of failed specimens of CFRP composites tested at 28 kg m/s. Figure 22 shows typical surfaces of failed specimens of unmodified CFRP composite, Figure 23 shows the same in the context of a CFRP composite modified with 0.5 wt. % SiC, while Figure 24 shows the same regarding a CFRP composite modified with 2.0 wt. % CS. The failure was typified by spliĴing the specimens into two, with the main failure mode being delamination. In addition, the failed specimens featured undulating, rough surfaces, which indicates that the matrix and fiber bundles ruptured during testing. This type of surface characteristic was more pronounced in SiC filler-modified CFRP composites than in those modified with CS filler. The surface morphology of com-

Figure 22. Typical optical macrographs of specimens of unmodified CFRP composites, based on HRJ-15881 resin, which were tested at 28 kg m/s. **Figure 22.** Typical optical macrographs of specimens of unmodified CFRP composites, based on HRJ-15881 resin, which were tested at 28 kg m/s. Polymers 2023, 15, x FOR PEER REVIEW 28 of 35

HRJ-15881 SP-6877

No No

Figure 23. Typical optical macrograph of a specimen of HRJ-15881-based CFRP composite modified with 0.5 wt.% SiC tested at 28 kg m/s. **Figure 23.** Typical optical macrograph of a specimen of HRJ-15881-based CFRP composite modified with 0.5 wt.% SiC tested at 28 kg m/s.

Figure 24. Typical optical macrograph a specimen of HRJ-15881-based CFRP composite modified with 2 wt.% CS tested at 28 kg m/s. **Figure 24.** Typical optical macrograph a specimen of HRJ-15881-based CFRP composite modified with 2 wt.% CS tested at 28 kg m/s.

between the matrix and CF could explain the reason why the fiber bundles remained in-

Figures 25–27 present SEM micrographs of specimens of modified and unmodified CFRP composites that were tested 28 kg m/s. The main failure modes of the unmodified CFRP composites (Figure 25) were CF bundle rupture, specimen spliĴing, and delamina-

tact despite rupture during impact testing.

Figures 25–27 present SEM micrographs of specimens of modified and unmodified CFRP composites that were tested 28 kg m/s. The main failure modes of the unmodified CFRP composites (Figure 25) were CF bundle rupture, specimen splitting, and delamination. Furthermore, in spite of the observed defects, it could be seen from Figure 25 that the CF fabric was impregnated by the phenolic resin quite well. The observed good adhesion between the matrix and CF could explain the reason why the fiber bundles remained intact despite rupture during impact testing. Polymers 2023, 15, x FOR PEER REVIEW 30 of 35

Figure 25. Typical SEM micrographs of unmodified HRJ-15881-based CFRP composites after SHPB testing at 28 kg m/s. **Figure 25.** Typical SEM micrographs of unmodified HRJ-15881-based CFRP composites after SHPB testing at 28 kg m/s.

Figure 26. Typical SEM fractographs of HRJ-15881-based CFRP composite modified with 0.5 wt.% SiC after SHPB testing at 28 kg m/s. **Figure 26.** Typical SEM fractographs of HRJ-15881-based CFRP composite modified with 0.5 wt.% SiC after SHPB testing at 28 kg m/s.

Figure 27 shows typical SEM micrographs of failed specimens of CFRP composites

Figure 27. Typical SEM fractographs of HRJ-15881-based CFRP composites modified with 2 wt.% CS after SHPB testing at 28 kg m/s. **Figure 27.** Typical SEM fractographs of HRJ-15881-based CFRP composites modified with 2 wt.% CS after SHPB testing at 28 kg m/s.

Plates of CFRP composites were produced using 2D woven CFs, microfillers, and two phenolic resins (SP-6877 and HRJ-15881). SHPB testing of specimens of the composites was conducted at two impact momentums, 15 kg m/s and 28 kg m/s. XRD diffractograms showed that the microfillers were intercalated in the phenolic resins. At 15 kg m/s, HRJ-15881-based CFRP composites modified with 1.5 wt.% SiC gave the highest impact strength. At 28 kg m/s, the impact strength of HRJ-15881-based CFRP composites was higher than that of the SP-6877-based CFRP composites at each level of SiC filler addition. In addition, in the range of SiC content used in this study (0.5–2.0 wt.%), the impact strength of SP-6877-based CFRP composites was beĴer than that of its composite containing no SiC filler. In contrast, only the impact strength of HRJ-15881-based composite modified with 0.5 wt.% SiC was higher than its composite without SiC filler. Nevertheless, the modification of the two resins with CS degraded the impact strength of the resulting com-

A microstructure examination showed that, at 15 kg m/s, no fracture occurred in the CFRP composites with and without microfillers. However, at 28 kg m/s, the CFRP composites failed by spliĴing, edge chip off, and crack extension through the matrix, which was more pronounced in the CFRP composites containing colloidal silica microfillers.

Author Contributions: Conceptualization, I.A.A., A.G.O. and I.N.A.O.; Methodology, I.A.A., A.G.O. and I.N.A.O.; Software, A.G.O.; Validation, I.A.A.; Formal analysis, I.A.A., O.E.A., I.N.A.O. and A.G.O.; Investigation, I.A.A.; Resources, I.N.A.O. and A.G.O.; Data curation, I.A.A. and O.E.A.; Writing—original draft preparation, I.A.A.; writing—review and editing, I.A.A., A.G.O. and I.N.A.O.; Visualization, I.A.A., O.E.A. and I.N.A.O.; Supervision, A.G.O. and I.N.A.O.; Project administration, I.A.A.; Funding acquisition, I.N.A.O. and A.G.O. All authors have read and agreed to

Funding: The financial support of the Natural Science and Engineering Research Council (NSERC) of Canada through the Discovery Grant award (RGPIN 2017-05751) is hereby acknowledged.

4. Conclusions

posites.

the published version of the manuscript.

Institutional Review Board Statement: Not applicable.

Informed Consent Statement: Not applicable.

Similarly, the failure mode of specimens of CFRP composites modified with 0.5 wt.% SiC (Figure 26) is also characterized by fiber bundle breakage and delamination. Nevertheless, a close comparison of Figures 25 and 26 shows that fiber bundle breakage is more pronounced in the SiC filler-modified composites than in the unmodified composites. This may explain the basis for the presence of undulating, jagged ridged surfaces in the optical images of tested specimen shown previously in Figures 22–24. Furthermore, Figure 26 shows there is fairly good bonding between the phenolic matrix and CF.

Figure 27 shows typical SEM micrographs of failed specimens of CFRP composites modified with 2 wt.% CS. As shown previously in Figures 25 and 26, the failure mode is typified by CF bundle breakage and delamination. A matrix crack is shown in the image on the left. Moreover, the micrographs indicate there is good bonding between the matrix and CF, which could be attributed to the good impregnation of CF preforms by the resins.

### **4. Conclusions**

The possibility of enhancing the resistance of CFRP composites to dynamic impact loading by modifying the matrix with SiC and CS microfillers was explored in this study. Plates of CFRP composites were produced using 2D woven CFs, microfillers, and two phenolic resins (SP-6877 and HRJ-15881). SHPB testing of specimens of the composites was conducted at two impact momentums, 15 kg m/s and 28 kg m/s. XRD diffractograms showed that the microfillers were intercalated in the phenolic resins. At 15 kg m/s, HRJ-15881-based CFRP composites modified with 1.5 wt.% SiC gave the highest impact strength. At 28 kg m/s, the impact strength of HRJ-15881-based CFRP composites was higher than that of the SP-6877-based CFRP composites at each level of SiC filler addition. In addition, in the range of SiC content used in this study (0.5–2.0 wt.%), the impact strength of SP-6877-based CFRP composites was better than that of its composite containing no SiC filler. In contrast, only the impact strength of HRJ-15881-based composite modified with 0.5 wt.% SiC was higher than its composite without SiC filler. Nevertheless, the modification of the two resins with CS degraded the impact strength of the resulting composites.

A microstructure examination showed that, at 15 kg m/s, no fracture occurred in the CFRP composites with and without microfillers. However, at 28 kg m/s, the CFRP composites failed by splitting, edge chip off, and crack extension through the matrix, which was more pronounced in the CFRP composites containing colloidal silica microfillers.

**Author Contributions:** Conceptualization, I.A.A., A.G.O. and I.N.A.O.; Methodology, I.A.A., A.G.O. and I.N.A.O.; Software, A.G.O.; Validation, I.A.A.; Formal analysis, I.A.A., O.E.A., I.N.A.O. and A.G.O.; Investigation, I.A.A.; Resources, I.N.A.O. and A.G.O.; Data curation, I.A.A. and O.E.A.; Writing—original draft preparation, I.A.A.; writing—review and editing, I.A.A., A.G.O. and I.N.A.O.; Visualization, I.A.A., O.E.A. and I.N.A.O.; Supervision, A.G.O. and I.N.A.O.; Project administration, I.A.A.; Funding acquisition, I.N.A.O. and A.G.O. All authors have read and agreed to the published version of the manuscript.

**Funding:** The financial support of the Natural Science and Engineering Research Council (NSERC) of Canada through the Discovery Grant award (RGPIN 2017-05751) is hereby acknowledged.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Acknowledgments:** The authors thank SI Group and Washington Mills for providing the phenolic resins and SiC microfiller used in this research, respectively, at no cost.

**Conflicts of Interest:** The authors performed the experimental work in this study independently. The authors declare that they have no known competing financial interest or personal relationships that could have appeared to influence the work reported in this paper.

### **References**

**Disclaimer/Publisher's Note:** The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

## *Article* **Research on Characterization of Nylon Composites Functional Material Filled with Al2O3 Particle**

**Jibing Chen 1,\* , Bowen Liu 1 , Maohui Hu 1 , Qianyu Shi 1 , Junsheng Chen 1 , Junsheng Yang 1 and Yiping Wu 2**

**\*** Correspondence<EMAIL_ADDRESS>
**Abstract:** This study revolves around the issues raised by the current semiconductor device metal casings (mainly composed of aluminum and its alloys), such as resource and energy consumption, complexity of the production process, and environmental pollution. To address these issues, researchers have proposed an eco-friendly and high-performance alternative material—Al2O3 particlefilled nylon composite functional material. This research conducted detailed characterization and analysis of the composite material through scanning electron microscopy (SEM) and differential scanning calorimetry (DSC). The results show that the Al2O3 particle-filled nylon composite material has a significantly superior thermal conductivity, about twice as high as that of pure nylon material. Meanwhile, the composite material has good thermal stability, maintaining its performance in high-temperature environments above 240 ◦C. This performance is attributed to the tight bonding interface between the Al2O3 particles and the nylon matrix, which not only improves the heat transfer efficiency but also significantly enhances the material's mechanical properties, with a strength of up to 53 MPa. This study is of great significance, aiming to provide a high-performance composite material that can alleviate resource consumption and environmental pollution issues, with excellent polishability, thermal conductivity, and moldability, which is expected to play a positive role in reducing resource consumption and environmental pollution problems. In terms of potential applications, Al2O3/PA6 composite material can be widely used in heat dissipation components for LED semiconductor lighting and other high-temperature heat dissipation components, thereby improving product performance and service life, reducing energy consumption and environmental burden, and laying a solid foundation for the development and application of future high-performance eco-friendly materials.

**Keywords:** PA6/Al2O3 composite; functional material; microstructure; mechanical property

Received: 15 April 2023 Revised: 15 May 2023 Accepted: 17 May 2023 Published: 19 May 2023

**Citation:** Chen, J.; Liu, B.; Hu, M.; Shi, Q.; Chen, J.; Yang, J.; Wu, Y. Research on Characterization of Nylon Composites Functional Material Filled with Al2O3 Particle. *Polymers* **2023**, *15*, 2369. https:// doi.org/10.3390/polym15102369 Academic Editor: S. D. Jacob Muthu

**Copyright:** © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

### **1. Introduction**

With the miniaturization, integration, high-density, and high-speed development of electronic components, heat dissipation has become a very serious problem [1], especially as various light-emitting diodes (LEDs) are widely used in various places for urban lighting. High-power and high-brightness LEDs generate a tremendous amount of heat, making it challenging to maintain the optimal operating temperature range of 25–30 ◦C for LEDs. Traditional inorganic thermally conductive materials such as metals and ceramics can no longer meet the needs of further development of electronic products, and new highefficiency thermally conductive materials must be used to dissipate heat promptly [2]. Polymer-based composite materials refer to the introduction of one or more inorganic or organic materials with specific properties into the polymer matrix through various processing methods (including chemical and physical means) so that the polymer also has some special properties, such as thermal conductivity, shielding of electromagnetic

waves, dielectric properties, electrical conductivity, and excellent damping performance [3]. Thermally conductive polymer composites have excellent thermal conductivity and unparalleled molding processability, electrical insulation, and corrosion resistance of traditional inorganic thermally conductive materials [4], becoming a research hotspot in the field of polymer materials.

The thermal conductivity of metals mainly relies on freely moving electrons, and the heat transfer mechanism is the formation of a heat conduction band using unconstrained electrons that can move freely; the thermal conductivity of inorganic non-metallic materials can only be conducted through the microscopic structure of mutual contact between material molecules by transferring phonons, using phonons as heat transfer carriers to transfer heat. Phonons are mechanical waves generated by lattice vibrations in crystals (mainly vibrations of atoms, molecules, and groups) [5,6], which are successively transferred between material molecules, so their thermal conductivity efficiency is not as good as that of free electrons. In contrast, polymers themselves are composed of molecular chains with very large molecular weights, and it is difficult to form a free electron energy band inside the material. Therefore, polymers mainly transfer heat through mutual contact between microscopic structures. However, unlike the phonon thermal conductivity of inorganic non-metallic materials, polymers themselves are composed of mixtures of homologous substances with different molecular weights, with larger van der Waals forces between molecular chains and longer chains that are prone to entanglement at the ends, resulting in more defects between the molecules. Therefore, polymer materials are particularly prone to phonon scattering, leading to low heat transfer efficiency and poor thermal conductivity of polymer materials [7]. Currently, there are two main approaches to preparing thermally conductive composite materials [8]: intrinsically thermally conductive polymer composites with inherent heat transfer capabilities, and filler-based thermally conductive polymer composites using externally added fillers. Intrinsic thermally conductive polymer materials have a large number of conjugated structures in their molecular chains and a relatively large proportion of crystalline regions, and these special structures give the polymer itself good heat transfer capabilities [9]. When phonons propagate in these polymers, there is less phonon scattering and hindrance, making intrinsic thermally conductive polymer composites excellent thermal conductors. However, the processing technology of intrinsic thermally conductive polymers is very complex, from the selection of initiators and the initiation of active polymer monomers, to the influence of reaction temperature and reaction time. These factors make the preparation of intrinsic thermally conductive polymers extremely difficult and the industrial production and processing difficult [10]. Therefore, based on the modern Fourier solid heat conduction theory, inorganic powders with strong heat transfer capabilities are filled into polymers to prepare high thermally conductive base polymer materials that meet the needs of the actual operating environment.

At present, a large number of studies have been conducted both domestically and abroad, aiming to develop composite functional materials with high thermal conductivity. Joao Paulo Berenguer, Arielle Berman, and other scholars [11] studied the effects of polyethylene fillers on the thermal conductivity, mechanical properties, and processability of all-polymer high thermal conductivity composite materials. Through thermal conductivity tests, tensile tests, bending tests, and impact tests, the results showed that optimizing the mass fraction and dispersion degree of polyethylene fillers can achieve a comprehensive optimization of the thermal conductivity, mechanical properties, and processability of all-polymer high thermal conductivity composite materials. Guorui Zhang, Sen Xue, and others [12] developed a high-efficiency thermal interface material with anisotropic orientation and high vertical thermal conductivity, by adjusting the distribution and orientation of graphene fillers, improving the thermal management performance of microelectronic devices, extending device life, and improving energy efficiency. You Y-L, Li D-X, and others [13] prepared GF/PA6 composite materials containing different solid lubricants in the laboratory, and used a ball-disc friction tester to evaluate the friction and wear performance of different composite materials, revealing the mechanism of various solid lubricants in

reducing friction and wear. One paper [14] mainly studied the thermal conductivity of polymer composite materials based on phenolic resin and boron nitride (BN), by adding boron nitride fillers to phenolic resin, improving the thermal conductivity of composite materials, and evaluating the thermal conductivity, thermal stability, and microstructure of composite materials through thermal conductivity tests, thermogravimetric analysis, and scanning electron microscopy. Another [15] analyzed the microstructure, thermal stability, and mechanical properties of Al2O3 and SiC-reinforced PA6 hybrid composite materials, and the experimental results showed that Al2O3 and SiC particles were uniformly distributed in the PA6 matrix, effectively improving the mechanical properties of the materials. Jiaqi Zhang, Xianzhao Jia, and others [16] mainly studied the effects of additive fluoroelastomer (FVMQ) and Al2O3 particle co-filling on its mechanical properties, thermal properties, and friction properties at high temperatures, verifying that by optimizing the mass fraction and dispersion degree of Al2O3 particle fillers, a comprehensive optimization of the mechanical properties, thermal properties, and friction properties of fluoroelastomer composite materials under high-temperature conditions can be achieved. Qiu Hong Mu, Dan Peng, and other scholars [17] studied the effect of filling Al2O3 particles on the thermal conductivity of silicone rubber, and evaluated the thermal conductivity, thermal stability, and microstructure of composite materials through thermal conductivity tests, thermogravimetric analysis, and scanning electron microscopy. One work [18] explored the effects of filling Al2O3 compounds on the thermal conductivity and rheological properties of epoxy resin and liquid crystal epoxy resin composite materials, by adding Al2O= compound fillers to epoxy resin and liquid crystal epoxy resin, improving the thermal conductivity and rheological properties of composite materials. Konopka K, Krasnowski M, and others [19] used pulsed plasma sintering (PPS) method to prepare Al2O3 samples and NiAl-Al2O3 composite materials, and analyzed the microstructure, phase composition, thermal stability, and mechanical properties of the two materials by X-ray diffraction (XRD), thermogravimetric analysis (TGA), and differential scanning calorimetry (DSC), and evaluated the significant influence of pulsed plasma sintering parameters on the microstructure and mechanical properties

According to the literature cited above, a series of research work has been carried out on thermal conductive materials for semiconductor devices. In order to ensure that these thermal conductive materials have excellent comprehensive performance, researchers usually choose to fill polymer materials with high thermal conductivity inorganic fillers or metal fillers. The thermal conductive materials prepared in this way have the advantages of low cost and easy processing. Through appropriate physical and chemical treatment methods or adjusting the experimental formula ratio, such thermal conductive materials can meet the specific application scenarios with thermal conductivity requirements. PA6 has good mechanical properties, thermal stability, and chemical stability, and can adapt to various complex working environments. More importantly, PA6 has high adjustability, and its performance can be improved by adding different types and proportions of fillers, reinforcements, and auxiliaries. Therefore, this study chose to add inorganic fillers (Al2O3) to nylon (PA6) and add 5% diethylhexyl phthalate (DEHP) in a certain proportion to prepare materials with specific functions. In order to study the mechanical properties of these functional materials, various instruments such as scanning electron microscopy (SEM) (JSM-7500F, JEOL Ltd., Tokyo, Japan), differential scanning calorimeter (DSC) (1/700 type, Mettler-Toledo, Zurich, Switzerland), and thermal conductivity meter (DTC-300, TA Instruments, New Castle, DE, USA) were used for testing and analysis. The results show that this high thermal conductivity functional material has great potential to replace metal shell materials and can be used to manufacture semiconductor devices with a metal appearance. This material performs excellently in polishing performance, thermal conductivity, and molding performance, providing promising possibilities for replacing metal materials. For example, in LED semiconductor lighting, it can be used as an effective heat dissipation material.

### **2. Materials and Methods**

### *2.1. Raw Materials*

The thermally conductive base material nylon (PA6) was purchased from Zhuhai SMIKA Polymer Company (Zhuhai, China); the thermally conductive filler Al2O3 (purity ≥ 99%) was obtained from National Medicine Group Chemical Reagent Co., Ltd. (Shanghai, China); the coupling agent diethylhexyl phthalate (DEHP) was provided by Foshan New Material Technology Co., Ltd. (Foshan, China).

### *2.2. Preparation of Materials*

Initially, the nylon plastic granules underwent essential pretreatment. The nylon granules were dried in an oven at a temperature of 120 ◦C for 24 h to eliminate moisture, ensuring the stability of material performance in subsequent experiments.

In this study, we conducted three sets of experiments. Firstly, to achieve a good thermal conductivity, thermal stability, and mechanical properties, we combined the research from ref. [20,21] and prior theoretical analysis. Nylon PA6, thermal conductive filler Al2O3, and DEHP were mixed at a mass ratio of 80:18:2. A high-speed stirrer was used to ensure that all components were fully mixed and formed a homogeneous mixture. Subsequently, the obtained mixture was extruded using an extruder to give the material a certain shape. The extruded material was then dried at 80 ◦C for 2 h to remove any residual moisture, ensuring the smooth progress of the subsequent injection molding process.

Secondly, with the aforementioned experimental procedure kept consistent, we conducted grouped experiments with filler particle size, filler shape, and filler ratio as variables to evaluate the thermal performance of the material. The particle sizes selected were 5 µm, 20 µm, and 100 µm; the filler shapes were spherical and flaky; the filler ratios were 10%, 20%, 30%, 40%, and 50%.

Lastly, to analyze the relationship between the filler ratio and mechanical properties in more detail, we set up nine groups of gradient experiments for the filler ratio, namely 0%, 5%, 10%, 15%, 20%, 25%, 30%, 35%, and 40%, while maintaining the aforementioned experimental procedure.

### *2.3. Characterization of Thermal Conductivity*

After the experiment was completed, we immersed the samples used for testing thermal performance in liquid nitrogen, allowing them to rapidly solidify and become brittle at low temperatures, thus facilitating subsequent cross-sectional observation and analysis. Then, the cross-section of the sample was placed on conductive adhesive and a sputter coating treatment performed to provide good conductivity and adhesion for observation and analysis under a scanning electron microscope (SEM). The SEM was used to observe the distribution characteristics and morphology of the thermally conductive filler in the matrix material. Finally, the prepared heat-resistant plastic samples were sent to a differential scanning calorimeter (DSC) for heat resistance performance analysis. We sought to understand the thermal stability and thermal degradation behavior of the composite material at different temperatures, thereby evaluating its application potential in hightemperature environments. Furthermore, we employed a thermal conductivity meter to analyze the samples. By applying heat to the sample and measuring its temperature change over a certain period, we calculated the sample's thermal conductivity, which allowed us to predict the material's thermal management performance in practical applications.

### *2.4. Mechanical Analysis*

In this study, tensile tests were conducted to evaluate the mechanical properties of the composite material, using an electronic universal testing machine (WDT-10, Shenzhen Kaiqiangli Experimental Instrument Co., Ltd., Shenzhen, China.) with a testing temperature of 25 ◦C. To ensure the accuracy and comparability of the experiments, the preparation and testing process of the tensile samples followed the relevant provisions of ASTM D638. Meanwhile, the size of the samples referred to the ISO 527-2 Type 1A standard requirements to ensure that the shape and dimensions of the samples met internationally accepted requirements. In the actual testing process, the samples were stretched at a transverse speed of 25 mm/min to simulate the loading speed and conditions that may be encountered in real-life applications. Finally, the American TA Instruments (New Castle, DE, USA) Q800 analyzer was used for single cantilever mode testing, gaining a deeper understanding of the mechanical properties of the composite material and its potential in various application scenarios. accepted requirements. In the actual testing process, the samples were stretched at a transverse speed of 25 mm/min to simulate the loading speed and conditions that may be encountered in real-life applications. Finally, the American TA Instruments (New Castle, DE, USA) Q800 analyzer was used for single cantilever mode testing, gaining a deeper understanding of the mechanical properties of the composite material and its potential in various application scenarios.

In this study, tensile tests were conducted to evaluate the mechanical properties of the composite material, using an electronic universal testing machine (WDT-10, Shenzhen Kaiqiangli Experimental Instrument Co., Ltd., Shenzhen, China.) with a testing temperature of 25 °C. To ensure the accuracy and comparability of the experiments, the preparation and testing process of the tensile samples followed the relevant provisions of ASTM D638. Meanwhile, the size of the samples referred to the ISO 527-2 Type 1A standard requirements to ensure that the shape and dimensions of the samples met internationally

### *2.5. Scanning Electron Microscopy (SEM) 2.5. Scanning Electron Microscopy (SEM)*  A scanning electron microscope (SEM) was used to study the microstructure of PA6,

A scanning electron microscope (SEM) was used to study the microstructure of PA6, Al2O3, and the composite material. Observing the microstructure and interactions within PA6, Al2O3, and their composite material helps to reveal the microstructural characteristics of the samples. Al2O3, and the composite material. Observing the microstructure and interactions within PA6, Al2O3, and their composite material helps to reveal the microstructural characteristics of the samples.

#### **3. Results and Discussion 3. Results and Discussion**

*2.4. Mechanical Analysis*

#### *3.1. Microstructure Analysis of Thermal Polymer Materials 3.1. Microstructure Analysis of Thermal Polymer Materials*

*Polymers* **2023**, *15*, x FOR PEER REVIEW 5 of 11

In order to measure the thermal conductivity, the test samples were first treated with liquid nitrogen to increase the brittleness of the material. Subsequently, the cross-section of the samples was placed on a conductive adhesive and subjected to sputter coating, enabling observation of the distribution characteristics and morphology of the thermally conductive fillers within the matrix material under a scanning electron microscope (SEM). Following this step, we utilized a differential scanning calorimeter (DSC) to analyze the thermal stability of the heat-resistant plastic samples to understand their performance under high-temperature conditions. Prior to the formal preparation of thermally conductive polymer materials, a detailed examination of the microstructure and energy spectrum distribution of the fillers and PA6 was conducted using a scanning electron microscope, based on the composition ratio of the thermally conductive fillers. The results are shown in Figures 1 and 2. In order to measure the thermal conductivity, the test samples were first treated with liquid nitrogen to increase the brittleness of the material. Subsequently, the cross-section of the samples was placed on a conductive adhesive and subjected to sputter coating, enabling observation of the distribution characteristics and morphology of the thermally conductive fillers within the matrix material under a scanning electron microscope (SEM). Following this step, we utilized a differential scanning calorimeter (DSC) to analyze the thermal stability of the heat-resistant plastic samples to understand their performance under high-temperature conditions. Prior to the formal preparation of thermally conductive polymer materials, a detailed examination of the microstructure and energy spectrum distribution of the fillers and PA6 was conducted using a scanning electron microscope, based on the composition ratio of the thermally conductive fillers. The results are shown in Figures 1 and 2.

**Figure 1. Figure 1.**  PA6 material microstructure image. PA6 material microstructure image.

After successfully preparing the thermally conductive composite material, its microstructure was observed using a scanning electron microscope (SEM), as shown in Figure 3. Comparing Figures 1 and 3, it can be found that the SEM image of the original PA6 displays a relatively uniform and smooth surface, while the SEM image of the PA6 composite material with added Al2O3 filler exhibits a clear distribution of filler particles and an interfacial region is formed between the Al2O3 filler and the PA6 matrix. Furthermore, the thermally conductive filler (Al2O3) appears as flake-like or spherical shapes and is relatively uniformly distributed within the nylon matrix material, which is beneficial for

forming continuous thermal conduction paths or networks within the composite material. After the addition of the thermally conductive filler, the original microstructure of the matrix material changes, resulting in the thermally conductive composite material exhibiting excellent thermal conductivity performance. *Polymers* **2023**, *15*, x FOR PEER REVIEW 6 of 11

*Polymers* **2023**, *15*, x FOR PEER REVIEW 6 of 11

**Figure 2.** Al2O3 energy spectrum analysis image. **Figure 2.** Al2O3 energy spectrum analysis image. iting excellent thermal conductivity performance.

**Figure 3.** SEM photographs of nylon composites: (**a**) 500 μm; (**b**) 100 μm. **Figure 3. Figure 3.** SEM photographs of nylon composites: (**a**) 500 µm; (**b**) 100 µm.

#### *3.2. DSC Testing 3.2. DSC Testing*

 (**a**) (**b**) **Figure 3.** SEM photographs of nylon composites: (**a**) 500 μm; (**b**) 100 μm. *3.2. DSC Testing*  In this study, the thermal properties of the thermally conductive polymer samples were analyzed using a differential scanning calorimeter (DSC). During the experiment, the samples were placed in a nitrogen atmosphere and heated at a rate of 20 °C per minute, In this study, the thermal properties of the thermally conductive polymer samples were analyzed using a differential scanning calorimeter (DSC). During the experiment, the samples were placed in a nitrogen atmosphere and heated at a rate of 20 °C per minute, with the temperature increasing from room temperature to 300 °C. After reaching 300 °C, the temperature was maintained at 10 °C for 10 min, followed by cooling. The obtained DSC curve of the samples is shown in Figure 4. Based on the results in Figure 4, the entire DSC curve can be divided into four stages [22]. In the first stage, from the start of heating to approximately 160 °C, the DSC curve gradually rises and reaches the peak of the first In this study, the thermal properties of the thermally conductive polymer samples were analyzed using a differential scanning calorimeter (DSC). During the experiment, the samples were placed in a nitrogen atmosphere and heated at a rate of 20 ◦C per minute, with the temperature increasing from room temperature to 300 ◦C. After reaching 300 ◦C, the temperature was maintained at 10 ◦C for 10 min, followed by cooling. The obtained DSC curve of the samples is shown in Figure 4. Based on the results in Figure 4, the entire DSC curve can be divided into four stages [22]. In the first stage, from the start of heating to approximately 160 ◦C, the DSC curve gradually rises and reaches the peak of the first endothermic peak; at this temperature, the polymer transitions from a glassy state to a highly elastic rubber-like or viscous fluid state, and the motion of the molecular chains becomes more flexible. The subsequent second stage extends from the end of the 160 ◦C endothermic peak to a temperature of 175 ◦C. The third stage has a temperature range of 175 ◦C to 240 ◦C, where the curve starts to rise, ultimately forming the second endothermic peak; at this temperature, the crystalline regions of PA6 begin to melt, and the movement of the molecular chains further increases. The final fourth stage covers a temperature range of 230 ◦C to 300 ◦C, during which the DSC curve begins to decline and exhibits a faster

with the temperature increasing from room temperature to 300 °C. After reaching 300 °C, the temperature was maintained at 10 °C for 10 min, followed by cooling. The obtained

DSC curve can be divided into four stages [22]. In the first stage, from the start of heating to approximately 160 °C, the DSC curve gradually rises and reaches the peak of the first

descent rate; in this process, heat is released, and the polymer material's chain segments begin to rearrange and form crystalline regions. ments begin to rearrange and form crystalline regions.

endothermic peak; at this temperature, the polymer transitions from a glassy state to a highly elastic rubber-like or viscous fluid state, and the motion of the molecular chains becomes more flexible. The subsequent second stage extends from the end of the 160 °C endothermic peak to a temperature of 175 °C. The third stage has a temperature range of 175 °C to 240 °C, where the curve starts to rise, ultimately forming the second endothermic peak; at this temperature, the crystalline regions of PA6 begin to melt, and the movement of the molecular chains further increases. The final fourth stage covers a temperature range of 230 °C to 300 °C, during which the DSC curve begins to decline and exhibits a faster descent rate; in this process, heat is released, and the polymer material's chain seg-

*Polymers* **2023**, *15*, x FOR PEER REVIEW 7 of 11

**Figure 4.** DSC curves of thermally conductive polymers. (Ⅰ) Glass transition phase; (Ⅱ) Thermal stabilization phase; (Ⅲ) Melting phase; (Ⅳ) Crystallization phase. **Figure 4.** DSC curves of thermally conductive polymers. (I) Glass transition phase; (II) Thermal stabilization phase; (III) Melting phase; (IV) Crystallization phase.

#### *3.3. Thermal Conductivity Testing 3.3. Thermal Conductivity Testing*

In this project, we conducted experiments on the thermal conductivity of polymer materials prepared based on Al2O3 with different diameters and morphologies. The diameters of the fillers used were 5 μm, 20 μm, and 100 μm, and the morphologies included spherical and flake-like shapes. Figure 5 shows the corresponding experimental data curves. Combining the experimental data with the literature [23], several conclusions can be drawn: (1) Under different diameters or different morphologies of the filler, the thermal conductivity of the polymer materials exhibits an increasing trend, indicating that the thermal conductivity of nylon composites increases with the increase in the content of Al2O3 filler. This phenomenon can be attributed to the increase in filler content, which helps form thermal conduction paths and networks in the composite material, thereby improving thermal conductivity; (2) By observing the data curves of the three groups with filler diameters of 5 μm, 20 μm, and 100 μm, it can be concluded that the thermal conductivity is the highest when the filler diameter is 20 μm and the lowest when the filler diameter is 100 μm. This indicates that within a specific particle size range, the thermal conductivity of the composite material increases with the increase in particle size; however, beyond this range, the thermal conductivity gradually decreases. This phenomenon may be due to the larger particle size filler reducing the thermal conductivity of the composite material, leading to a decrease in overall thermal conductivity; (3) Comparing the data of spherical and flake-like fillers with a diameter of 5 μm, it can be found that the flake-like fillers are more effective in improving the thermal conductivity of the composite material than spherical fillers. In this project, we conducted experiments on the thermal conductivity of polymer materials prepared based on Al2O3 with different diameters and morphologies. The diameters of the fillers used were 5 µm, 20 µm, and 100 µm, and the morphologies included spherical and flake-like shapes. Figure 5 shows the corresponding experimental data curves. Combining the experimental data with the literature [23], several conclusions can be drawn: (1) Under different diameters or different morphologies of the filler, the thermal conductivity of the polymer materials exhibits an increasing trend, indicating that the thermal conductivity of nylon composites increases with the increase in the content of Al2O3 filler. This phenomenon can be attributed to the increase in filler content, which helps form thermal conduction paths and networks in the composite material, thereby improving thermal conductivity; (2) By observing the data curves of the three groups with filler diameters of 5 µm, 20 µm, and 100 µm, it can be concluded that the thermal conductivity is the highest when the filler diameter is 20 µm and the lowest when the filler diameter is 100 µm. This indicates that within a specific particle size range, the thermal conductivity of the composite material increases with the increase in particle size; however, beyond this range, the thermal conductivity gradually decreases. This phenomenon may be due to the larger particle size filler reducing the thermal conductivity of the composite material, leading to a decrease in overall thermal conductivity; (3) Comparing the data of spherical and flake-like fillers with a diameter of 5 µm, it can be found that the flake-like fillers are more effective in improving the thermal conductivity of the composite material than spherical fillers. *Polymers* **2023**, *15*, x FOR PEER REVIEW 8 of 11

*3.4. Mechanical Properties Testing*

various filler percentages.

even lower than that of single nylon material.

**Figure 5.** The thermal conductivity of spherical and flaky Al2O3 filled with nylon composites with 5, 20, and 100 μm particles. **Figure 5.** The thermal conductivity of spherical and flaky Al2O3 filled with nylon composites with 5, 20, and 100 µm particles.

Figure 6 shows the change curve of the tensile strength of composite materials as the

of the composite material reaches its maximum value of 55.1 MPa, which is about 19% higher than that of pure nylon material. However, when the alumina filler content exceeds 14%, the tensile strength of the composite material significantly decreases, showing an inverse relationship with the proportion of filler. When the alumina content reaches 30%, the tensile strength is 35.5 MPa. This observation indicates that the tensile strength of the composite material increases with the increase in alumina content within a specific range (alumina content below 10%). When the content is below this range, its tensile strength is

**Figure 6.** Tensile strength curves of alumina/nylon composites. ■ denotes the tensile strength at

With the addition of alumina filler, the compressive strength of nylon composite materials has been improved accordingly. An electronic universal testing machine was used to measure the compressive strength of composite materials, and the experimental results are shown in Figure 7. When the alumina content reaches 15%, the compressive strength

#### *3.4. Mechanical Properties Testing* Figure 6 shows the change curve of the tensile strength of composite materials as the

*3.4. Mechanical Properties Testing*

5, 20, and 100 μm particles.

Figure 6 shows the change curve of the tensile strength of composite materials as the content of Al2O3 filler increases. According to the experimental data curve, the tensile strength of nylon is 46 MPa. When the alumina content reaches 14%, the tensile strength of the composite material reaches its maximum value of 55.1 MPa, which is about 19% higher than that of pure nylon material. However, when the alumina filler content exceeds 14%, the tensile strength of the composite material significantly decreases, showing an inverse relationship with the proportion of filler. When the alumina content reaches 30%, the tensile strength is 35.5 MPa. This observation indicates that the tensile strength of the composite material increases with the increase in alumina content within a specific range (alumina content below 10%). When the content is below this range, its tensile strength is even lower than that of single nylon material. content of Al2O3 filler increases. According to the experimental data curve, the tensile strength of nylon is 46 MPa. When the alumina content reaches 14%, the tensile strength of the composite material reaches its maximum value of 55.1 MPa, which is about 19% higher than that of pure nylon material. However, when the alumina filler content exceeds 14%, the tensile strength of the composite material significantly decreases, showing an inverse relationship with the proportion of filler. When the alumina content reaches 30%, the tensile strength is 35.5 MPa. This observation indicates that the tensile strength of the composite material increases with the increase in alumina content within a specific range (alumina content below 10%). When the content is below this range, its tensile strength is even lower than that of single nylon material.

**Figure 5.** The thermal conductivity of spherical and flaky Al2O3 filled with nylon composites with

*Polymers* **2023**, *15*, x FOR PEER REVIEW 8 of 11

**Figure 6.** Tensile strength curves of alumina/nylon composites. ■ denotes the tensile strength at various filler percentages. **Figure 6.** Tensile strength curves of alumina/nylon composites. denotes the tensile strength at various filler percentages.

With the addition of alumina filler, the compressive strength of nylon composite materials has been improved accordingly. An electronic universal testing machine was used to measure the compressive strength of composite materials, and the experimental results are shown in Figure 7. When the alumina content reaches 15%, the compressive strength With the addition of alumina filler, the compressive strength of nylon composite materials has been improved accordingly. An electronic universal testing machine was used to measure the compressive strength of composite materials, and the experimental results are shown in Figure 7. When the alumina content reaches 15%, the compressive strength of the composite material reaches its highest value of 94 MPa. Compared with pure PA6, the compressive strength of this composite material has increased by 47.6%. Within a certain range (alumina content ≤ 15%), the compressive strength of nylon composite materials increases with the increase in alumina content. This trend can be attributed to the dispersion of alumina particles in the polymer matrix and the interaction between alumina particles and the polymer matrix [24]. However, when the alumina content exceeds this threshold, the excessive filler content leads to particle aggregation in the polymer matrix, which affects the overall performance of the composite material, resulting in a decreasing trend in compressive strength.

trend in compressive strength.

**Figure 7.** Compressive strength curve of Al2O3/nylon composites. ▲ denotes the compressive strength at various filler percentages. **Figure 7.** Compressive strength curve of Al2O3/nylon composites. N denotes the compressive strength at various filler percentages.

of the composite material reaches its highest value of 94 MPa. Compared with pure PA6, the compressive strength of this composite material has increased by 47.6%. Within a certain range (alumina content ≤ 15%), the compressive strength of nylon composite materials increases with the increase in alumina content. This trend can be attributed to the dispersion of alumina particles in the polymer matrix and the interaction between alumina particles and the polymer matrix [24]. However, when the alumina content exceeds this threshold, the excessive filler content leads to particle aggregation in the polymer matrix, which affects the overall performance of the composite material, resulting in a decreasing

#### **4. Conclusions 4. Conclusions**

In this paper, researchers successfully prepared a composite material with high thermal conductivity using high thermal conductive functional materials, namely polyamide 6 (PA6) and alumina (Al2O3). Detailed experimental studies were conducted on the thermal conductivity, heat resistance, and performance at different filler ratios of this composite material. The experimental results show that: (1) using spherical fillers with a diameter of 20 μm can achieve a larger thermal conductivity. When the filler size is reduced to 5 μm, the thermal conductivity reaches the highest value of 0.9 W·(mK)−1. In addition, this composite material has good heat resistance performance and can maintain stability at the highest temperature of 240 °C. (2) The composite material has significant mechanical properties. When the Al2O3 content is 10%, the tensile strength of the thermal conductive functional material reaches its highest value. However, when the additive amount is increased to 14%, the compressive strength of the composite material reaches the highest value of 55.1 MPa, indicating that to a certain extent, increasing the Al2O3 content can improve the mechanical properties of the composite material. In summary, this high thermal conductive functional material prepared from PA6 and Al2O3 not only has excellent thermal conductivity and heat resistance performance but also has good mechanical properties. This gives this composite material a wide range of application potential in hightemperature application fields, such as aerospace, automotive manufacturing, and electronic devices. This research provides beneficial practical experience for engineers and scientists in related fields, helping to develop higher performance composite materials in In this paper, researchers successfully prepared a composite material with high thermal conductivity using high thermal conductive functional materials, namely polyamide 6 (PA6) and alumina (Al2O3). Detailed experimental studies were conducted on the thermal conductivity, heat resistance, and performance at different filler ratios of this composite material. The experimental results show that: (1) using spherical fillers with a diameter of 20 µm can achieve a larger thermal conductivity. When the filler size is reduced to 5 µm, the thermal conductivity reaches the highest value of 0.9 W·(mK)−1 . In addition, this composite material has good heat resistance performance and can maintain stability at the highest temperature of 240 ◦C. (2) The composite material has significant mechanical properties. When the Al2O3 content is 10%, the tensile strength of the thermal conductive functional material reaches its highest value. However, when the additive amount is increased to 14%, the compressive strength of the composite material reaches the highest value of 55.1 MPa, indicating that to a certain extent, increasing the Al2O3 content can improve the mechanical properties of the composite material. In summary, this high thermal conductive functional material prepared from PA6 and Al2O3 not only has excellent thermal conductivity and heat resistance performance but also has good mechanical properties. This gives this composite material a wide range of application potential in high-temperature application fields, such as aerospace, automotive manufacturing, and electronic devices. This research provides beneficial practical experience for engineers and scientists in related fields, helping to develop higher performance composite materials in actual applications.

actual applications. **Author Contributions:** Conceptualization, J.C. (Jibing Chen) and B.L.; methodology, J.C. (Jibing Chen); software, J.C. (Jibing Chen) and B.L.; validation, J.C. (Junsheng Chen), B.L. and M.H.; formal **Author Contributions:** Conceptualization, J.C. (Jibing Chen) and B.L.; methodology, J.C. (Jibing Chen); software, J.C. (Jibing Chen) and B.L.; validation, J.C. (Junsheng Chen), B.L. and M.H.; formal analysis, J.Y. and Y.W.; investigation, J.C. (Junsheng Chen) and B.L.; resources, J.C. (Junsheng Chen); data curation, M.H. and Q.S.; writing—original draft preparation, B.L., J.C. (Junsheng Chen) and M.H.; writing—review and editing, Q.S. and Y.W.; supervision, J.C. (Jibing Chen), Y.W. and J.Y.; project administration, J.C. (Jibing Chen) and Y.W.; funding acquisition, J.C. (Jibing Chen) and J.Y. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was supported by the Open Fund of Hubei Longzhong Laboratory (2022ZZ-17) and the Science and Technology Project of Science and Technology Department of Hubei Province (No. 2022EHB020).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Acknowledgments:** The authors acknowledge the technical support of Hubei Longzhong Laboratory, and the Analytical and Testing Center of HUST for their analytical work.

**Conflicts of Interest:** The authors declare no conflict of interest.

### **References**

**Disclaimer/Publisher's Note:** The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

## *Article* **Morphological, Mechanical, and Antimicrobial Properties of PBAT/Poly(methyl methacrylate-***co***-maleic anhydride)–SiO2 Composite Films for Food Packaging Applications**

**Raja Venkatesan 1,† , Krishnapandi Alagumalai 1,†, Chaitany Jayprakash Raorane 1 , Vinit Raj 1 , Divya Shastri 2,\* and Seong-Cheol Kim 1,\***

**\*** Correspondence<EMAIL_ADDRESS>(D.S<EMAIL_ADDRESS>(S.-C.K.); Tel.: +82-53-810-2787 (S.-C.K.)

† These authors contributed equally to this work.

**Abstract:** A poly(methyl methacrylate-*co*-maleic anhydride) P(MMA-*co*-MA) copolymer was synthesized via radical polymerization. The synthesized P(MMA-*co*-MA) copolymer was identified by 1Hand 13C-nuclear magnetic resonance spectroscopy (1H-NMR), (13C-NMR), Fourier-transform infrared (FTIR) spectroscopy, X-ray diffraction (XRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). The poly(butylene adipate-*co*-terephthalate) (PBAT)/P(MMA-*co*-MA)– SiO2 composites were developed using a solution-casting method. The PBAT to P(MMA-*co*-MA) weight ratio was kept at 70:30, while the weight percentage of SiO2 nanoparticles (NPs) was varied from 0.0 to 5.0 wt.%. SiO2 was used for PBAT/P(MMA-*co*-MA) to solve the compatibility between PBAT and the P(MMA-*co*-MA) matrix. The PBAT/P(MMA-*co*-MA)–SiO2 composites were characterized by studied FTIR spectroscopy, XRD, SEM, and TEM. A comparison of the composite film PBAT/P(MMA-*co*-MA)–SiO2 (PBMS-3) with the virgin PBAT and P(MMA-*co*-MA) film revealed its good tensile strength (19.81 MPa). The WVTR and OTR for the PBAT/P(MMA-*co*-MA)–SiO2 composites were much smaller than for PBAT/P(MMA-*co*-MA). The PBAT/P(MMA-*co*-MA)–SiO2 WVTR and OTR values of the composites were 318.9 ± 2.0 (cc m−2 per 24 h) and 26.3 ± 2.5 (g m−2 per 24 h). The hydrophobicity of the PBAT/P(MMA-*co*-MA) blend and PBAT/P(MMA-*co*-MA)–SiO2 composites was strengthened by the introduction of SiO2 , as measured by the water contact angle. The PBAT/P(MMA-*co*-MA)–SiO2 composite films showed excellent antimicrobial activity against the food-pathogenic bacteria *E. coli* and *S. aureus* from the area of inhibition. Overall, the improved packaging characteristics, such as flexibility, tensile strength, low O2 and H2O transmission rate, and good antimicrobial activities, give the PBAT/P(MMA-*co*-MA)–SiO2 composite film potential for use in food packaging applications.

**Keywords:** PBAT; P(MMA-*co*-MA) copolymer; SiO2 nanoparticles; PBAT/P(MMA-*co*-MA)–SiO2 composites; mechanical strength; antimicrobial activity; food packaging

### **1. Introduction**

The advances in food packaging materials in recent decades have attracted considerable attention because of the increasing need for effectively packaged foods with an extended shelf life [1]. The active film can offer extra safety features in addition to the usual barrier characteristics [2,3]. The antioxidant and antimicrobial activities are significant because they help alleviate concerns with food quality, such as degradation, rancidity, and coloration. An active food packaging material should fulfill a list of needs, including having sufficient physical and mechanical properties, a high barrier, a high level of barrier performance, safe for human health, and no harmful impact on the environment [4]. Furthermore, a straightforward, reproducible, and economical preparation method is essential for large-scale applications [5,6]. New production techniques that enhance the safety of

**Citation:** Venkatesan, R.; Alagumalai, K.; Raorane, C.J.; Raj, V.; Shastri, D.; Kim, S.-C. Morphological, Mechanical, and Antimicrobial Properties of PBAT/Poly(methyl methacrylate-*co*-maleic anhydride)–SiO2 Composite Films for Food Packaging Applications. *Polymers* **2023**, *15*, 101. https:// doi.org/10.3390/polym15010101

Academic Editor: S. D. Jacob Muthu

Received: 17 November 2022 Revised: 17 December 2022 Accepted: 21 December 2022 Published: 26 December 2022

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

food packaging are available. The manufacture of food packaging must ensure that the packaging volume and weight are kept to a minimum to maintain the required stability, hygienic conditions, and consumer acceptance. However, there are significant obstacles to using biodegradable food packaging in markets. Technologies that enhance the oxygen and water barrier of biodegradable polymer systems are ideal for biodegradable packaging in daily life. In recent times, notable potential improvements have been adopted, showing promise in fabricating high oxygen/water vapor-permeable biodegradable materials for food packaging.

Radical polymerization is the primary method for synthesizing poly(methyl methacrylate*co*-maleic anhydride) P(MMA-*co*-MA) [7]. Pre-polymerization and cast polymerization are the two steps that are always involved in polymerization [8]. Furthermore, P(MMA*co*-MA) copolymer can be used in ring-opening reactions to form cube-shaped materials that can be used in various applications, such as electrolytes for batteries, surfactants and emulsified solutions, and antiwear additives [9,10]. It is often used to fabricate composite materials that transfer heat easily because of its excellent properties [11–13]. P(MMA*co*-MA) is the focus of the current research, which shows a relationship with particle size [14], flame retardants [15], inorganic nanoparticles [16], bio-additives [17], chemical modification [18], and polymerization conditions [19]. On the other hand, the brittleness and poor mechanical characteristics of P(MMA-*co*-MA) constitute a significant impediment to its broad application. Some research has been conducted to develop P(MMA-*co*-MA) composites with reinforcement materials, such as polymers, chitin nanofibers, and cellulose, to increase the mechanical and thermal properties and solve the issues [11]. Nevertheless, there are still a few disadvantages. For example, it cannot substantially enhance the mechanical properties of P(MMA-*co*-MA) materials, and using nanofillers can increase the cost of P(MMA-*co*-MA) composites, which is another impediment to scaling up production. Impurities cause opacity and the use is confined because of the hydrolysis reaction of the maleic anhydride groups [20,21]. Maleic anhydride (MA) and methyl methacrylate (MMA) were copolymerized. Other components may alter P(MMA-*co*-MA) to circumvent these limitations [12,22].

As a transparent thermoplastic, lightweight, or shatter-resistant replacement for glassware, poly(butylene adipate-*co*-terephthalate) (PBAT) was blended with P(MMA-*co*-MA). It is non-toxic and contains a strong hydrogen bonding character. It can also be used as a casting polymer in inks and coatings. Owing to its plasticizing property [9], PBAT increases the flexibility and barrier characteristics of the produced composite film, which are essential for packaging applications [23,24]. The most notable biocompatibility polymers in the thermoplastic materials group, PBAT, are synthesized via reactions among the monomers adipic acid, terephthalic acid, and 1,4-butanediol [25]. This material is attractive because of its excellent physical features, high elongation at break, and amorphous structure. In addition, it can be obtained using eco-friendly and renewable technologies [26,27]. The capability to produce scaffolding structures enables PBAT to be used in tissue engineering to reconstruct or replace tissue (such as bone, cartilage, and blood vessels). On the other hand, reinforcing fillers must strengthen the PBAT matrix's performance and enhance its mechanical and thermal properties for broader use. Countless loadings in the PBAT matrices have been studied to improve the physical, thermal, permeability, and characteristics of the finished goods.

Nano-silica is used in combination with PBAT/P(MMA-*co*-MA) blends as a nanofiller to improve the characteristics of materials owing to its low cost, high reproducibility, and suitability for large-scale production [28]. SiO2 nanoparticle (NP) synthesis is a topic of interest in basic and applied research. Other variables include the particle-size suitability for multiple products. The last few years have seen a significant increase in interest in SiO2 filled composites. According to studies, composite materials loaded with SiO2 possess good thermal and mechanical characteristics [29,30]. Studies have been carried out on the SiO2 distribution in poly(methyl methacrylate) matrices [31,32], high-density polyethylene [33], and poly(ethylene oxide) [34]. PBAT is a thermoplastic compostable material commonly

used to manufacture films, bottles, and fibers for different applications because of its outstanding mechanical and optical characteristics, susceptibility to fatigue and wear, and creep-fragmentation resistance [35,36]. Functional characteristics, such as antimicrobial activity, have been enhanced during the fabrication of PBAT-based composite films [37–39].

Recent research has been conducted on biomass or sustainable polymers reinforced with fillers and fabricated to use the solution casting method [40–42]. In this work, a series of PBAT/P(MMA-*co*-MA) composites with different contents was prepared by solution casting of PBAT/P(MMA-*co*-MA) blends and a SiO2 nanopowder [43]. SiO2 has a disk structure according to scanning electron microscopy (SEM). The excellent compatibility and interactions between SiO2 and PBAT mean that SiO2 NPs are dispersed uniformly in PBAT/P(MMA-*co*-MA) blends according to the shape of the film surface [44]. The tensile strength of the PBAT/P(MMA-*co*-MA) composites increased from 9.43 MPa for PBAT without reinforcement to 22.82 MPa with 5.0 wt.% of SiO2 NPs (PBMS-3). Thermogravimetric analysis (TGA) of the composite films of PBAT/P(MMA-*co*-MA)–SiO2 showed an increase in the weight loss temperature (%). The PBAT/P(MMA-*co*-MA)–SiO2 composite materials also have high hydrophobicity. The antimicrobial activity of SiO2-incorporated PBAT/P(MMA-*co*-MA) composites is remarkable because of the combined antimicrobial activity of SiO2 NPs against *S. aureus* and *E. coli*. As a result, composites composed of PBAT/P(MMA-*co*-MA)–SiO2 were developed and assessed for food packaging applications.

### **2. Materials and Methods**

### *2.1. Chemicals and Materials*

Poly(butylene adipate-*co*-terephthalate) (PBAT) pellets were supplied by BASF. The melt flow index (MFI) (190 ◦C, 2.16 kg) was 3.3–6.6 g/10 min. A solution of methyl methacrylate (MMA) was purchased from Daejung Chemicals in Korea. Maleic anhydride (MA; 98%, Sigma-Aldrich, St. Louis, MO, USA) was recrystallized from chloroform and n-hexane. By crystallizing from methanol, the radical initiator of azobisisobutyronitrile (AIBN; 99%, Sigma–Aldrich, St. Louis, MO, USA) was purified. Tetraethoxysilane (TEOS) was received from Sigma–Aldrich in India. Daejung Chemicals in Korea supplied methyl benzene, ethyl acetate, ethanol, and N, N-dimethylformamide (DMF). All compounds were used as received. For the preparation of solutions and reactions, double-distilled water was used.

### *2.2. Synthesis of SiO2 Nanoparticles (NPs)*

SiO2 NPs were synthesized using the methodology reported elsewhere [45]. Tetraethoxysilane (TEOS) was used as a silica source in the sol–gel technique to prepare SiO2 NPs. The reaction was rapidly stirred for 30 min at 150 ◦C with 7.4 mL TEOS and 80 mL methanol introduced. A white precipitate was heated for one hour at 350 ◦C in a tubular furnace to generate SiO2. The XRD, SEM, and TEM characterization data matched the information from the previous report [46].

### *2.3. Synthesis of Poly(methyl methacrylate-co-maleic anhydride)*