Research works [7,8] show that continuous nanofibers have a higher molecular orientation with a better degree of crystallinity and, thus, they improve the properties of the interlaminar regions. Wang et al. [9] cited that continuous nanofibers could improve the mechanical properties along a particular direction and promote load-bearing capabilities. These results expounded the innovative ideas in manufacturing aligned and continuous nanofibers in the form of nanomats. These nanomats could develop a bridging mechanism within the interlaminar regions and be suitable interlaminar reinforcements for glass fiber hybrid composites. The continuous-aligned nanofibers (CANFs) mats will be thin, light in weight, and will not increase the weight of final composites. Moreover, they can reduce the length scale mismatch between the macroscale fibers and the polymer matrix molecular chains. In addition, these nanomats could be collected directly over the fiber mats and thus reduce the processing time/cost.

Several studies focus on using various nanoparticles prepared as mats for reinforcing composites such as graphene-embedded paper for electromagnetic interference shielding (EMI) [10],graphene/SiCnw nanostrctured films for improving the mechanical and thermal properties of carbon fiber composites [11]. Among them, the single-walled carbon nanotubes (SWCNTs), though they are comparatively expensive, have attracted researchers due to their exceptional mechanical properties with higher specific surface area [12]. However, the inherent issues, such as agglomeration, alignment, and hydrophobic nature, limit their potential as secondary reinforcements, which necessitated the innovative research ideas for using SWCNTs as interlaminar reinforcements [13]. One such direction could be to develop continuous-aligned nanofibers mats doped with SWCNTs.

Several techniques are being used for producing polymeric nanofibers, and among them, electrospinning is considered simple and versatile. However, the polymeric jets formed during electrospinning have two distinct phases: near field, and far field [14]. In the near field, the jet takes a straight path, and in the far field, it starts whipping into a complex path due to electrical, gravitational, and rheological forces. Though the far field yields nanosized fibers [15], the whipping leads to randomly oriented nanofibers, which defeats the objective of producing continuous and aligned nanofibers. Moreover, the nanosized fibers are vital for aligning and dispersing SWCNTs into the nanofibers mats. To overcome the whipping instability, researchers have used thin static collectors [16] or high-speed rotating collectors [17]. However, based on our experience, a wide rotating drum collector with auxiliary vertical electrodes (AVEs) seems to be the most suitable solution for controlling the whipping effect and also for improving the alignment of nanofibers within the nanomats. Moreover, AVEs have better control over the whipping instability at the far field, as confirmed by other researchers [18].

Hence, the focus of this study is to utilize the electrospinning setup with AVEs for producing CANFs mats and functionalized SWCNTs (f-SWCNTs) doped CANFs mats. The SWCNTs will be functionalized to improve the interfacial adhesion with nanofibers. Lastly, the effect of CANFs and f-SWCNTs doped CANFs mats on the mechanical properties of glass fiber composites will be evaluated using tensile, flexural, and impact tests to characterize the interlaminar strengthening mechanism.

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

### *2.1. Materials*

The polymer solution used in electrospinning was prepared by mixing Polyacrylonitrile powder (PAN: density 1.184 g/cm3 at 25 ◦C) with N, N-dimethylformamide (DMF: 99% purity) supplied by Sigma Aldrich, Kempton Park, South Africa. The solution was stirred for 22 h at 50 ◦C until it became homogeneous. The composites were fabricated using Bisphenol-A epoxy (AMPREG 21 Resin) with woven E-glass mats supplied by AMT

Composites Pty. Ltd., Johannesburg, South Africa. The SWCNTs with 99% purity were used to dope CANFs for strengthening the interlaminar region. The SWCNTs and the required functionalization chemicals, such as Polyvinyl alcohol (PVA) and dimethyl sulfoxide (DMSO), were purchased from Sigma Aldrich, Kempton Park, South Africa. The physical and mechanical properties of Bisphenol-A epoxy, woven E-glass, and SWCNTs are given in Table 1.

**Table 1.** Properties of materials used—data collected from the suppliers.

### *2.2. Testing Procedure*

The tensile, three-point flexural, and Charpy impact tests were carried out on the composite specimens as per ASTM D638-08: 2010, ASTM D256-06a: 2010, and ASTM D790-07: 2010, respectively. The specimen thickness was kept at 3.2 mm. The flexural test specimens were fabricated with the thickness to width ratio of 1:16, and the gauge length between the supports was kept at 63 mm. A 20 kN SHIMADZU tensile testing machine was used to conduct both tensile and flexural tests (Figure 1a,b, which were carried out at a cross-head speed of 2 mm/min. The Charpy impact tests were carried out on a Tinius Olsen impact pendulum as per the ISO 179-1 standard (Figure 1c. The breaking energy was used to determine the impact resistance of the composite specimens.

**Figure 1.** (**a**) Tensile testing setup and specimens. (**b**) Flexural testing setup and specimens. (**c**) Impact testing setup and specimens.

### *2.3. Morphological Characterization*

Morphological characterization of nanofibers and hybrid composite specimens was carried out using scanning electron microscopy (SEM-Nova 600 Nanolab and Quanta 400 FEG from FEI Company, Hillsboro, OR, USA) at the voltage range of 5 KV to 30 KV. Before analysis, the specimens were coated with a thin layer of carbon and a 15 nm layer of gold–palladium (Au/Pd) in a ratio of (60:40) in an EMITECH K950X evaporator and EMITECH K550X sputter coater, respectively. The specimens were then attached to 1 cm diameter stubs with DAG 580 colloidal graphite. The FTIR analysis was conducted on a Bruker Vector 22 FTIR spectrometer using 64 sample scans and 32 background scans. The Raman analysis was conducted using a Horiba Jobin-Yvon Lab RAM HR Raman spectrometer with a 514.5 nm line of an argon-ion laser as the excitation source.

### *2.4. SWCNTs Functionalization*

The SWCNTs were functionalized using the Friedel–Crafts alkylation (FCA) process. The FCA process was carried out by mixing SWCNTs (0.4 g) and PVA (4 g) with a 40 mL DMSO solution in a 250 mL flask using a magnetic stirrer. The flask was then immersed in an oil bath for gradual heating to 90 ◦C in the presence of nitrogen. As the heating promotes the chemical process, a catalyst of aluminum chloride (ALCl3) was added for grafting the PVA chains onto the surface of the SWCNTs. During the chemical process, the OH groups from PVA reacted with ALCl3 and produced positive carbocations, (CH3+ ions : electrophiles), and negative ALCl− 2 ions, respectively. The electrophiles were then attracted to the SWCNTs' hexagonal ring due to the CNT's surface delocalized electron (nucleophilic) sites, and thus attached the PVA chains onto the SWCNTs. After 20 h, the chemical reaction was aborted by adding 100 mL of methanol/hydrochloric acid mixture (volumetric ratio 1:1). The solution was then centrifuged at 3500 rpm for 15 min to precipitate the functionalized SWCNTs (f- SWCNTs). The f-SWCNTs were washed, filtered, and dried in a furnace for 3 h at 70◦ .

The Raman spectra were used to analyze the functionalization results of both pristine (p-SWCNTs) and functionalized SWCNTs (f-SWCNTs) and are shown in Figure 2i. The two peaks at the wavelengths of 1300–1400 cm−1 and 1500–1600 cm−1 were for the disorder (D band) and graphite band (G band), respectively. These bands represent two different vibrational modes of SWCNTs. The D band represents the crystal disorder, such as sidewall defects (pentagons or heptagons) or Sp3 carbon hybridization, and the G band was for the Sp2 carbon hybridization. The ratio of D band to G band (*ID*/*IG*) increases with the increasing disorder [19] and defines the structural changes of the SWCNTs wall due to functionalization [20]. In the current work, the *ID*/*IG* ratio was measured at 0.1025 and 0.1674 for p-SWCNTs and f-SWCNTs, respectively. The higher ratio of 0.1674 for f-SWCNTs was due to the carbocations attachment to the SWCNTs walls and the conversion of Sp2 to Sp3 hybridized carbons. This confirmed the PVA grafting to the SWCNTs wall during the FCA process [21].

Figure 2ii shows the FTIR spectrum for the range of 4000 to 500 cm−1 of both p-SWCNTs and f-SWCNTs. By analyzing the f-SWCNTs spectrum, the peak at 1028 cm−1 showed the C–O stretching, while the peak from 1100 to 1600 cm−1 confirmed the aromatic structures. The peaks corresponding to –CH and –OH bonds were obtained at 2900 cm−1 and 3200–3400 cm−1 , respectively. These peaks confirm that the PVA chains are attached to the surface of SWCNTs.

**Figure 2.** (**i**) Raman spectra of p-SWCNTs and f-SWCNTs. (**ii**) FTIR spectra of (**a**) p-SWCNTs and (**b**) f-SWCNTs.

### *2.5. Electrospinning of Aligned and Doped Nanofibers Mats*

The electrospinning setup considered for this research consists of a DC voltage supply, NE-1600 syringe pump, and a grounded rotating drum collector (diameter 120 mm and width 210 mm), respectively. The electrospun nanomats were collected over the glass fiber mats attached to the rotating drum collector. The details of the complete setup with the spinning parameters were presented in our published manuscript [22] and the readers are encouraged to refer to the paper for further details. Following our previous experimental results, three electrospinning parameters, i.e., applied voltage, the distance between electrodes, and PAN concentration, were investigated in this work. In addition, auxiliary vertical electrodes (AVEs) were placed between the spinneret and collector. The range of tested and the optimized values of the spinning parameters are given in Table 2, which produced the nanofiber mats with fiber diameters in the range of 120–150 nm.

**Table 2.** Spinning parameter values with AVEs.

To analyze the effect of AVEs on the whipping effect, an electric field simulation software was used to simulate the electrostatic forces. The software developed the electrostatic forces as electric field lines (green lines) between the spinneret and the collector. The positive electrode (the spinneret) is red, while the blue color refers to the negative electrode. The simulation shows that without AVEs, the electric field lines are spread over the entire width of the collector. As the electric fields were not focused on the collector, the whipping effect deposited the nanofibers randomly over the whole width of the collector (Figure 3a). However, after introducing AVEs (Figure 3b), the simulation shows that the electric field lines were narrowed down towards the center of the collector (nanofiber driving zone) by the AVEs electric field. This phenomenon reduced the whipping effect and aligned the nanofibers perpendicular to the collector axis. Figure 4 shows the camera images and scanning electron microscope (SEM) images of electrospun nanofibers mats produced without (Figure 4a) and with (Figure 4b) AVEs. As it can be seen, the AVEs significantly improved the alignment of nanofibers and further reduced the diameter of the nanofibers to 150 nm, which is 1/3 smaller than the nanofibers produced without AVEs. This confirms that the addition of AVEs significantly refined the nanofiber morphologies.

**Figure 3.** Electric field simulation software analysis (**a**) without AVEs; (**b**) with AVEs.

The electrospinning setup was used to produce the required nanomats. At first, the randomly oriented PAN nanofiber (RONFs) mats were produced without AVEs for the weight fractions of 0.1%, 0.2%, 0.5%, and 1%, respectively. Based on the experimental results, the weight fraction of 0.5 wt% PAN nanofibers were selected to manufacture continuousaligned nanofibers (CANFs) mats using AVEs. Finally, the 0.5 wt% CANFs mats were

doped with 0.25 wt% f-SWCNTs to produce f-SWCNTs doped CANFs mats. For doping, only 0.25 wt% f-SWCNTs were selected to minimize the density effect following the results from Pilehrood et al. [23]. Before spinning the nanomats, the required weight fractions of PAN-DMF solution with and without f-SWCNTs were prepared and electrospun over the glass fiber mats.

## (b) Aligned Nanofibres Mat

**Figure 4.** Electrospun nanofibers mats (**a**) without AVEs; (**b**) with AVEs.

### *2.6. Composites Manufacturing Process*

The composites were fabricated using a vacuum-assisted resin transfer molding (VARTM) method. Following our previous work [4], 32 vol% glass fiber composites were selected as a reference for producing nanomats strengthened hybrid composites. Considering the manuscript length, the details of the glass fiber composites experimental results are not included, and the mechanical properties for the 32 vol% glass fiber composites are given in Table 3.

The VARTM mold consists of a female square steel ring, and upper and lower male plates, and these were assembled with the alternative layers of woven E-glass fiber mats coated with the RONFs, CANFs, and f-SWCNTs doped CANFs mats. Then, the resinhardener mixer was placed at the mold suction port and the vacuum pressure of 80 psi was maintained until the resin filled the mold. Finally, the resin was allowed to cure for 24 h at room temperature. Figure 5 shows the manufacturing process of nanomats strengthened

hybrid composites specimens, which were machined into tensile, flexural, and impact test specimens.

**Table 3.** Properties of 32 vol% glass fiber composites.

**Figure 5.** Fabrication process of glass fiber hybrid composites preform.

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

The discussions of the results are divided into two sections. The first section will focus on analyzing the results from RONFs mats strengthened hybrid composite specimens, and the CANFs mats and f-SWCNTs doped CANFs mats strengthened hybrid composites will be discussed in the second section.

### *3.1. RONFs Mats Strengthened Hybrid Composites*

Figure 6 shows the mechanical properties of RONFs mats strengthened hybrid composites, and it uses the 32 vol% glass fiber composites as reference. The initial addition of 0.1 wt% of mats showed a scanty increase in properties compared to the 32 vol% glass fiber composites. Further addition of 0.5 wt% RONF mats showed an increase in tensile strength from 156.52 to 251.51 MPa, elastic modulus from 11.77 to 14.36 GPa, flexural strength from 242.2 to 404.01 MPa, flexural modulus from 9.58 GPa to 14.58 GPa, and impact resistance from 160.18 KJ/m2 to 193.112 KJ/m2 , respectively. The statistical parameters for the RONFs strengthened glass hybrid composites are presented in Table 4.

The increase in the properties could be attributed to the reinforcing effect of RONFs mats due to their higher specific surface area, which could have increased the contact area within interlaminar regions. The infused RONFs mats provided more resistance to the relative slippage of the matrix over the glass fiber mats while transferring the applied load (refer to Figure 7a) and thus improved the mechanical properties. In addition, the geometrical web network of RONFs mats provided bridging between the glass fibers and the matrix, as shown in Figure 7b, and thus increased the resistance to the applied load and improved the properties.

**Table 4.** Statistical parameters of the RONFs strengthened glass fiber composites.

**Figure 6.** Properties of RONFs mats strengthened hybrid composite: (**a**) tensile properties; (**b**) flexural properties; (**c**) impact properties.

However, a similar effect was not observed for further increase in RONFs mats from 0.5 wt% to 1 wt%. This may be associated with the nanofibers agglomeration and poor penetration of resin into the nanofibers network due to the random orientation with higher weight fractions, which could have acted as stress concentration sites. Figure 8 provides evidence of agglomerated nanofibers within the interlaminar regions.

### *3.2. CANFs and f-SWCNTs Doped CANFs Mats Strengthened Hybrid Composites*

Based on the above results, 0.5 wt% of nanomats were selected to manufacture 0.5 wt% CANFs mats and 0.25 wt% f-SWCNTs doped with 0.5 wt% CANFs mats strengthened glass fiber hybrid composites specimens to understand the effect of alignment and f-SWCNTs doping within the interlaminar regions. The experimental results used 0.5 wt% of RONF mats strengthened hybrid composites as a reference, and the results are presented below.

**Figure 7.** Fracture behavior of 0.5 wt% RONFs mats strengthened hybrid composite (**a**) Pulled out nanofibres (**b**) Bridging of nanofibres.

**Figure 8.** Agglomerated nanofibers of 1 wt% RONFs mats strengthened hybrid composite.

### Mechanical Properties

The tensile, flexural, and impact properties of CANFs mats and f-SWCNTs doped CANFs mats strengthened hybrid composites are shown in Figure 9, and the properties values with percentage of increase are shown in Table 5.

**Figure 9.** Mechanical properties of 0.5 wt% CANFs and 0.25 wt% f-SWCNTs doped with 0.5 wt% CANFs mats strengthened hybrid composites': (**a**) tensile properties; (**b**) flexural properties; (**c**) impact properties.

**Table 5.** RONF mats , CANF mats, and f-SWCNTS doped CANF mats hybrid composites properties and their increase.

The phenomenal increase in mechanical properties was possible due to the reinforcement effect of continuous/aligned nanofibers along with uniformly dispersed f-SWCNTs doped CANFs mats within the interlaminar regions. The increase in mechanical properties can be attributed to the molecular orientations, reduction in nanofibers diameter, and the uniform dispersion of the aligned nanofibers network due to AVEs. Further, the reduction in nanofibers diameters (450 nm to 150 nm) provided an approximately threefold increase in nanofibers network within the interlaminar region.

Figure 10 shows the fractured surfaces of RONFs, CANFs, and f-SWCNTs doped CANFs mats strengthened hybrid composites. These images confirm that the RONFs mats (Figure 10a) accumulated nanofibers and created bigger voids as they were pulled out from the matrix. This led to an inefficient strengthening mechanism due to which the initiated

cracks propagated without much resistance and resulted in poor mechanical properties. On the other hand, the CANFs mats with reduced diameter and improved distribution and orientation showed an orderly pull out of nanofibers and broken heads of nanofibers within the matrix (Figure 10c). This indicates that the bridging mechanism created by the nanomats networks acted as barriers for the propagating cracks, and, as such, the interlaminar regions were strengthened by the CANFs mats.

**Figure 10.** Fracture surfaces: (**a**) Agglomeration of RONFs mats strengthened hybrid composite tensile test specimen. (**b**) Tensile test specimen fracture surface of f-SWCNTs doped CANFs mats strengthened hybrid composite. (**c**) Fracture surface of the interlaminar regions with aligned nanomats networks of flexural test specimen. (**d**) Branched cracks within the interlaminar region due to nanofibers network.

Figure 10b shows the SEM images of the fractured surface of f-SWCNTs doped CANFs mats strengthened hybrid composite. The images show several propagating cracks around the f-SWCNTs doped nanofibers heads. This implies that the initiated cracks were arrested by the f-SWCNTs doped CANFs mats and resisted/diverted the propagating cracks. The resistance prevented the separation of the matrix due to the propagating cracks. This phenomenon is evident in Figure 10d, where the cracks were either stopped or diverted due to the presence of f-SWCNTs doped nanofibers (refer to sites 1, 2, and 3). At site 3, the crack traveled across the nanofiber after the f-SWCNTs doped nanofiber was pulled out, and in sites 1 and 2, the cracks propagated over the broken nanofiber heads or pulled out f-SWCNTs doped nanofibers. This confirmed that the increase in mechanical properties of hybrid composites is due to the resistive f-SWCNTs doped nanofibers mats network within the interlaminar regions.

As discussed before, the bridging mechanism within the multiscale hybrid composites can be seen from the SEM images, as shown in Figure 11a,b. The functionalization of SWCNTs provided improved the dispersion, and further strengthened the bonding between f-SWCNTs and nanofibers, which could have resulted in a stronger nanofibers network within the interlaminar regions. Further, the alignment provided the required bridging, which resisted the crack propagation and promoted load transfer within the interlaminar

region. Thus, the f-SWCNTs doped aligned nanomats strengthened hybrid composite provided the highest mechanical properties.

**Figure 11.** Fracture surfaces of 0.5 wt% nanofibers strengthened hybrid composites (**a**) Elongated f-SWCNTs doped CANFs mats in bridging. (**b**) Bridging of f-SWCNTs doped nanofibers. (**c**) Fractured surfaces of impact tested specimen. (**d**) Pulled and broken head of f-SWCNTs.

### **4. Conclusions and Recommendations**

The continuously aligned nanofiber mats with and without f-SWCNTs were produced using an electrospinning setup with AVEs, which were then used to reinforce the interlaminar regions of glass fiber composites. The effect of aligned nanomats on the mechanical properties of glass fiber composites was seen from the experimental results; as such, both CANFs and f-SWCNTs doped CANFs produced hybrid composites with improved properties by strengthening the interlaminar properties. The addition of AVEs with the electrospinning reduced the nanofiber diameter and improved their alignments within the interlaminar regions. The reduced diameter and alignment of nanofibers helped to improve the molecular orientation and further helped to distribute f-SWCNTs uniformly into the glass fiber hybrid composites. The functionalization of SWCNTs promoted better dispersion of SWCNTs and improved the bonding between f-SWCNTs and nanofibers. These factors promoted better load transfer within the interlaminar regions and thus improved the mechanical properties of the glass fiber composites. The research could serve in different applications, such as electronics, filtration, etc., where aligned nanofibers with reduced diameter and improved molecular orientation are the essential requirements. The application of SWCNTs could also be extended to other research areas, such as smart textiles, transistors, information storage devices, batteries, EMI shielding, etc.

**Author Contributions:** Conceptualization and investigation: A.M. and M.N. Experiments and resources: A.M. and N.A. Writing—review and editing: A.M., N.A. and J.M. Supervision and funding acquisition: J.M. All authors have read and agreed to the published version of the manuscript.

**Funding:** This project was partially supported by the funding from Natural Sciences and Engineering Research Council of Canada (DDG-2020-00046) and National Research Foundation South Africa.

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

**Data Availability Statement:** Data are contained within the article.

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

### **References**

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## *Article* **Effects on the Thermo-Mechanical and Interfacial Performance of Newly Developed PI-Sized Carbon Fiber–Polyether Ether Ketone Composites: Experiments and Molecular Dynamics Simulations**

**Hana Jung † , Kwak Jin Bae † , Yuna Oh, Jeong-Un Jin , Nam-Ho You and Jaesang Yu \***

> Composite Materials Application Research Center, Institute of Advanced Composite Materials, Korea Institute of Science and Technology, Chudong-ro 92, Bongdong-eup, Wanju-gun 55324, Republic of Korea

**\*** Correspondence<EMAIL_ADDRESS>
† These authors contributed equally to this work.

**Abstract:** In this study, polyether ether ketone (PEEK) composites reinforced with newly developed water-dispersible polyimide (PI)-sized carbon fibers (CFs) were developed to enhance the effects of the interfacial interaction between PI-sized CFs and a PEEK polymer on their thermo-mechanical properties. The PI sizing layers on these CFs may be induced to interact vigorously with the p-phenylene groups of PEEK polymer chains because of increased electron affinity. Therefore, these PI-sized CFs are effective for improving the interfacial adhesion of PEEK composites. PEEK composites were reinforced with C-CFs, de-CFs, and PI-sized CFs. The PI-sized CFs were prepared by spin-coating a water-dispersible PAS suspension onto the de-CFs, followed by heat treatment for imidization. The composites were cured using a compression molding machine at a constant temperature and pressure. Atomic force and scanning electron microscopy observations of the structures and morphologies of the carbon fiber surfaces verified the improvement of their thermomechanical properties. Molecular dynamics simulations were used to investigate the effects of PI sizing agents on the stronger interfacial interaction energy between the PI-sized CFs and the PEEK polymer. These results suggest that optimal amounts of PI sizing agents increased the interfacial properties between the CFs and the PEEK polymer.

**Keywords:** carbon fiber; polyimide; sizing agent; interfacial adhesion; thermo-mechanical properties; molecular dynamics simulations

### **1. Introduction**

In recent years, polyether ether ketone (PEEK) polymer composites reinforced with carbon fibers (CFs) have been widely used in high-performance applications, such as in the aerospace and automobile fields, due to their high heat resistance, very light weight, and great mechanical strength [1–14]. However, PEEK polymer has a high melting temperature of around 343 ◦C, which is constrained during the manufacturing process to improve the physical properties of the resulting composites [15–19]. Therefore, it was necessary to develop a sizing agent to protect the CFs with better thermal stability to increase the interfacial performance between the carbon fibers and matrix, which is an important factor in the load transfer of composites.

A variety of treatment methods (e.g., plasma, synthesis, and inducing functional groups of sizing agents on the fiber surfaces) have been studied [20–29]. Naito [30] demonstrated an increase in the tensile properties of composites reinforced with carbon nanotube (CNT)-grafted and polyimide (PI)-coated CFs. Hassan et al. [31] proposed that a loosely packed CNT network using PI could improve the interfacial adhesion between CFs and matrix due to enhanced CF wettability, polarity, and roughness. As a compatibilizer, PI

**Citation:** Jung, H.; Bae, K.J.; Oh, Y.; Jin, J.-U.; You, N.-H.; Yu, J. Effects on the Thermo-Mechanical and Interfacial Performance of Newly Developed PI-Sized Carbon Fiber–Polyether Ether Ketone Composites: Experiments and Molecular Dynamics Simulations. *Polymers* **2023**, *15*, 1646. https:// doi.org/10.3390/polym15071646

Academic Editor: S. D. Jacob Muthu

Received: 27 February 2023 Revised: 21 March 2023 Accepted: 23 March 2023 Published: 25 March 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/).

exhibits a relatively high molecular weight in the interface regions as a result of a synergetic effect between PI and CNT. However, because PI is hard to dissolve, it requires a large amount of a strong organic solvent (such as N-Methyl-2-pyrrolidone) and limits the enhancement of interfacial adhesion [32–35]. Therefore, there is a need for novel research to create PI with high heat resistance for use as a sizing agent for carbon fibers.

Molecular dynamics (MD) simulations are attractive for predicting the effects of surface characteristics on the physical properties of polymer-based composites reinforced with carbon material [36–38]. It is possible to examine the thermal behavior of polymer-based composites by calculating the movement of molecules over time, which cannot be confirmed via experiments. Jiao et al. [39] and Jin et al. [40] studied the interfacial properties between polymers and sized-, functionalized-carbon materials using MD simulations. However, these results are not considered to prove enhanced interfacial properties without molecular mobility, which is significant in the thermal behavior of composites.

In this study, the thermo-mechanical properties and optimal sizing agent content of CF/PEEK composites with sizing agents, which can be applied even at the high processing temperature of PEEK polymers, were investigated. The aqueous PI sizing agent effects on the structure, morphology, interface performance, and thermo-mechanical properties of CF/PEEK composites were analyzed in detail to prove the comprehensive contributions of environmentally friendly modifications using various surface analysis techniques (Fouriertransform infrared spectroscopy (FT-IR), nuclear magnetic resonance (NMR), atomic force microscopy (AFM), and field-emission scanning electron microscopy (FE-SEM)) and physical estimates (interlaminar shear strength (ILSS), and dynamic mechanical analyzer (DMA) and thermogravimetric analyzers (TGA) characteristics). MD simulations were performed to analyze the interfacial characterization and thermal behavior according to the effect of the PI sizing agent between the CFs and the PEEK polymer. The calculated results can validate the experiments of the CF/PEEK composites.

### **2. Experiments**

### *2.1. Materials*

Pyromellitic dianhydride (PMDA, >99%), 4,40 -oxydianiline (ODA, >98%), and 4-hydroxy-1-methylpiperidine (HMP, >98%) were purchased from Tokyo Chemical Industry (Tokyo, Japan) to synthesize the poly (amic acid) salt (PAS). The N-Methyl-2 pyrrolidinone (NMP) was obtained from Sigma-Aldrich (St. Louis, MO, USA). A CF plain-weave fabric (C120-3K, plain, HD FIBER Co., Ltd., Yangsan, Republic of Korea) was used as reinforcement for the composites. C120-3K has a density of 1.78 g·cm−3 , a weight of 203 g·m−2 , and a thickness of 0.25 mm. The amorphous PEEK film (APTIV 2000, Victrex, Co., Ltd., Lancashire, UK) was used as a composite matrix with high heat resistance. The PEEK film has a density of 1.26 g·cm−3 , a tensile modulus of 1.80 GPa, a thickness of 16 µm, and a glass transition and melting temperature of 143 and 343 ◦C, respectively.

### *2.2. Preparation of Poly (Amic Acid) Salt (PAS) as a PI Sizing Agent*

The polyimide (PI) was prepared from poly (amic acid) (PAA), a precursor, using a two-step reaction. The first step was the formation of the PAA solution via a nucleophilic substitution reaction of the monomer in a step-growth polymerization reaction. The second step was the dehydration and cyclization of the synthesized PAA solution with thermal treatment over 200 ◦C resulting in the final PI sizing agent. The synthetic process of the PAS powder was as follows (Figure 1). An amount of 0.15 mol of ODA was dissolved in 1200 mL of NMP at room temperature. Then, 0.15 mol of PMDA was added to the ODA solution, and the resulting mixture was stirred for 24 h under N2 gas. This resulted in the formation of a PAA solution. Then, 0.3 mol of HMP was added to the PAA solution and stirred for 1 h. The mixture was then precipitated in acetone, and the resulting product was filtered, washed, and dried under vacuum at room temperature to obtain a PAS powder [15].

der [15].

**Figure 1.** Synthetic route for PAS. **Figure 1.** Synthetic route for PAS.

#### *2.3. Fabrication of Water−Dispersible PI−Sized CF/PEEK Composites 2.3. Fabrication of Water-Dispersible PI-Sized CF/PEEK Composites*

Suspensions were created to deploy the sizing agent onto the carbon fiber surfaces using the PAS powder in distilled water. The water−dispersible PAS powder was heated in a water bath at 60 °C for 4 h to obtain the PI sizing agent, which was homogenized on a hot plate by stirring it with a magnetic bar. Sizing agents were prepared with PI contents of 0.5, 1.0, and 1.5 wt% compared with the CFs. The commercial carbon fabric (C−CF) was immersed in acetone at room temperature for 2 h to remove epoxy groups from the sizing agent. Then, the carbon fabric was cleaned in acetone in 2 steps (10 min each) and dried in a vacuum oven at 40 °C. The de−sized carbon fabric (de−CF) was modified by coating it with the prepared PAS suspensions using a spin coater at 800 rpm for 30 sec. The PAS−modified CFs with different contents of 0.5, 1.0, and 1.5 wt% were dried in an oven at 40 °C for 12 h. For the imidization of the PAS−modified CF, heat treatment was carried out in an oven at 250 °C for 1 h. An illustration of the surface treatment of the CFs with the PI sizing agent is shown in Figure 2. In this study, the laminate was composed of 25 layers of carbon fiber fabric and 150 plies of PEEK film. PEEK composites were reinforced with C−CF, de−CF, and PI−sized CF. The laminate was preheated using a compression−molding machine at a constant temperature of 380 °C for 10 min without pressure. After the preheating process, the temperature was maintained at 380 °C with a pressure of 10 MPa for 20 min. The cooling step was performed using a water−cooling system with a rate of -10 °C∙min-1 . The total weight fraction of the carbon fibers was about 58 wt%. Suspensions were created to deploy the sizing agent onto the carbon fiber surfaces using the PAS powder in distilled water. The water-dispersible PAS powder was heated in a water bath at 60 ◦C for 4 h to obtain the PI sizing agent, which was homogenized on a hot plate by stirring it with a magnetic bar. Sizing agents were prepared with PI contents of 0.5, 1.0, and 1.5 wt% compared with the CFs. The commercial carbon fabric (C-CF) was immersed in acetone at room temperature for 2 h to remove epoxy groups from the sizing agent. Then, the carbon fabric was cleaned in acetone in 2 steps (10 min each) and dried in a vacuum oven at 40 ◦C. The de-sized carbon fabric (de-CF) was modified by coating it with the prepared PAS suspensions using a spin coater at 800 rpm for 30 s. The PAS-modified CFs with different contents of 0.5, 1.0, and 1.5 wt% were dried in an oven at 40 ◦C for 12 h. For the imidization of the PAS-modified CF, heat treatment was carried out in an oven at 250 ◦C for 1 h. An illustration of the surface treatment of the CFs with the PI sizing agent is shown in Figure 2. In this study, the laminate was composed of 25 layers of carbon fiber fabric and 150 plies of PEEK film. PEEK composites were reinforced with C-CF, de-CF, and PI-sized CF. The laminate was preheated using a compression-molding machine at a constant temperature of 380 ◦C for 10 min without pressure. After the preheating process, the temperature was maintained at 380 ◦C with a pressure of 10 MPa for 20 min. The cooling step was performed using a water-cooling system with a rate of −10 ◦C·min−1 . The total weight fraction of the carbon fibers was about 58 wt%.

was filtered, washed, and dried under vacuum at room temperature to obtain a PAS pow-

**Figure 2.** Schematic diagram of the process for surface treatment of carbon fibers using the PAS solution.

#### **Figure 2.** Schematic diagram of the process for surface treatment of carbon fibers using the PAS solution. *2.4. Characterizations*

*2.4. Characterizations* Nuclear magnetic resonance (NMR) spectra were determined with an Agilent 600 MHz Premium COMPACT spectrometer at 600 MHz for 1H in dimethyl sulfoxide−*d*6 (DMSO−*d*6), using tetramethylsilane (TMS) as an internal standard. The attenuated total reflection–Fourier−transform infrared (ATR−FT−IR) spectra were obtained with a Nicolet IS10 with 32 scans per spectrum at a 2 cm–1 resolution. Atomic force microscopy (AFM) was conducted with a peak force tapping mode system (multimode−8, BRUKER Co., Billerica, MA, USA) to measure the surface roughness of the carbon fibers. Field−emission scanning electron microscopy (FE−SEM Verios 460L, FEI Corp., Hillsboro, OR, USA) with energy−dispersive X−ray spectroscopy (EDX) was performed to characterize the surface morphologies of the carbon fibers as well as the fracture surfaces of the composites reinforced with carbon fibers. The interlaminar shear strength (ILSS) of the composites was measured using the short−beam shear test as three−point bending according to the ASTM D 2344 standard (5569, Instron, Norwood, MA, USA) [41]. Specimens with the dimensions 28.2 mm × 9.4 mm × 4.7 mm were tested at a cross−head speed rate of 1 mm∙min−1 . The short−beam strength (*F*sbs) was calculated from the equation *F*sbs = 0.75 Pb/*bh*, where Pb represents the maximum load, b is the width, and h is the thickness of the specimen. The sample results were estimated as an average value from more than five specimens. The thermo−mechanical properties of the composites were obtained using a dynamic mechanical analyzer (DMA Q800, TA Instruments, New Castle, DE, USA) with a dual−cantilever mode at 1 Hz from 30 °C to 300 °C and a heating rate of 3 °C∙min−1 , according to the ASTM 4065−01 standard [42]. The thermal stability of the carbon fibers was analyzed using thermogravimetric−differential thermal analysis (TG−DTA Analyzer SDT Q600, WATERS, Nuclear magnetic resonance (NMR) spectra were determined with an Agilent 600 MHz Premium COMPACT spectrometer at 600 MHz for 1H in dimethyl sulfoxide-*d*6 (DMSO-*d*6), using tetramethylsilane (TMS) as an internal standard. The attenuated total reflection– Fourier-transform infrared (ATR-FT-IR) spectra were obtained with a Nicolet IS10 with 32 scans per spectrum at a 2 cm−1 resolution. Atomic force microscopy (AFM) was conducted with a peak force tapping mode system (multimode-8, BRUKER Co., Billerica, MA, USA) to measure the surface roughness of the carbon fibers. Field-emission scanning electron microscopy (FE-SEM Verios 460L, FEI Corp., Hillsboro, OR, USA) with energy-dispersive X-ray spectroscopy (EDX) was performed to characterize the surface morphologies of the carbon fibers as well as the fracture surfaces of the composites reinforced with carbon fibers. The interlaminar shear strength (ILSS) of the composites was measured using the short-beam shear test as three-point bending according to the ASTM D 2344 standard (5569, Instron, Norwood, MA, USA) [41]. Specimens with the dimensions 28.2 mm × 9.4 mm × 4.7 mm were tested at a cross-head speed rate of 1 mm·min−1 . The short-beam strength (*F* sbs) was calculated from the equation *F* sbs = 0.75 Pb/*bh*, where Pb represents the maximum load, b is the width, and h is the thickness of the specimen. The sample results were estimated as an average value from more than five specimens. The thermo-mechanical properties of the composites were obtained using a dynamic mechanical analyzer (DMA Q800, TA Instruments, New Castle, DE, USA) with a dual-cantilever mode at 1 Hz from 30 ◦C to 300 ◦C and a heating rate of 3 ◦C·min−1 , according to the ASTM 4065-01 standard [42]. The thermal stability of the carbon fibers was analyzed using thermogravimetric-differential thermal analysis (TG-DTA Analyzer SDT Q600, WATERS, Milford, MA, USA) with a temperature range of 25 ◦C to 800 ◦C with a heating rate of 10 ◦C min−1 under nitrogen gas.

#### Milford, MA, USA) with a temperature range of 25 °C to 800 °C with a heating rate of 10 *2.5. Molecular Dynamics Modeling*

°C min-1 under nitrogen gas. *2.5. Molecular Dynamics Modeling* MD simulations were performed to investigate the effect of the PI sizing agent on the interface between the CFs and PEEK polymer using Materials Studio 2017 software. The simulation model consisted of five graphene sheets, PI molecular chains, and PEEK molecular chains for modeling (Figure 3). The PI molecular chains with 0.0, 0.3, 1.0, 2.0, and 5.0 wt% compared with the carbon fibers were placed on the top surfaces of the graphene sheets. The simulation of the system based on a constant number, volume, and temperature (NVT) ensemble was performed to equilibrate all the models with a time step of 1.0 fs for 300 ps at 300 K. The simulation of the system based on a constant number, pressure, and temperature (NPT) ensemble was performed to further relax the internal stresses of the simulation models with an additional annealing process. All the models were heated from 300 K to 700 K at 50 K intervals and cooled down to 300 K at equal intervals. Each MD simulations were performed to investigate the effect of the PI sizing agent on the interface between the CFs and PEEK polymer using Materials Studio 2017 software. The simulation model consisted of five graphene sheets, PI molecular chains, and PEEK molecular chains for modeling (Figure 3). The PI molecular chains with 0.0, 0.3, 1.0, 2.0, and 5.0 wt% compared with the carbon fibers were placed on the top surfaces of the graphene sheets. The simulation of the system based on a constant number, volume, and temperature (NVT) ensemble was performed to equilibrate all the models with a time step of 1.0 fs for 300 ps at 300 K. The simulation of the system based on a constant number, pressure, and temperature (NPT) ensemble was performed to further relax the internal stresses of the simulation models with an additional annealing process. All the models were heated from 300 K to 700 K at 50 K intervals and cooled down to 300 K at equal intervals. Each step was performed with a time step of 1.0 fs during 300 ps. The density of the equilibrated CF/PEEK models was 1.54 g·cm−3 , which was similar to the experimental result of 1.55 g·cm−3 . The dimension of the CF/PEEK composite model was 48.2 Å (*x*) × 49.9 Å (*y*) × 53.1 Å (*z*) in the

periodic boundary condition. All the MD simulations were performed using the condensedphase-optimized molecular potentials for the atomistic simulation studies (COMPASS II) force field [43], which has been widely used to describe the interaction between carbon materials and polymers. Å (*z*) in the periodic boundary condition. All the MD simulations were performed using the condensed−phase−optimized molecular potentials for the atomistic simulation studies (COMPASS ΙΙ) force field [43], which has been widely used to describe the interaction between carbon materials and polymers.

. The dimension of the CF/PEEK composite model was 48.2 Å (*x*) × 49.9 Å (*y*) × 53.1

step was performed with a time step of 1.0 fs during 300 ps. The density of the equilibrated

, which was similar to the experimental result of 1.55

*Polymers* **2023**, *15*, 1646 5 of 17

CF/PEEK models was 1.54 g∙cm-3

g∙cm-3

**Figure 3.** Molecular structures and composite models in the MD simulations. **Figure 3.** Molecular structures and composite models in the MD simulations.

The mean square displacement (MSD) can be used to characterize the speed at which particles move and quantify the mobility of molecules according to temperature. The MSD could be determined in the temperature range from 250 K to 500 K during the NPT ensemble. The MSD analysis can be calculated using the following equation: The mean square displacement (MSD) can be used to characterize the speed at which particles move and quantify the mobility of molecules according to temperature. The MSD could be determined in the temperature range from 250 K to 500 K during the NPT ensemble. The MSD analysis can be calculated using the following equation:

$$\text{MSD} = \frac{1}{N} \sum\_{i=1}^{N} \left| \mathbf{x}^{(i)}(t) - \mathbf{x}^{(i)}(0) \right|^2 \tag{1}$$

where is the number of particles to be averaged, () (0) is the reference position of the −th particle, and () () is the position of the −th particle at time . where *N* is the number of particles to be averaged, *x* (*i*) (0) is the reference position of the *i*-th particle, and *x* (*i*) (*t*) is the position of the *i*-th particle at time *t*.

#### *2.6. Interfacial Characterization 2.6. Interfacial Characterization*

The interaction energy of the CFs with the PEEK polymer was calculated using The interaction energy of the CFs with the PEEK polymer was calculated using

$$
\Delta E \text{ (kcal/mol)} = E\_{\text{total}} - \left( E\_{\text{CF}} + E\_{\text{polymer}} \right) \tag{2}
$$

mer, represents to the total potential energy of the composite, and and are the potential energy of the PI−sized CFs and PEEK polymer, respectively. The interfacial shear strength (ISS, *τ*) between the PI−sized CFs and the PEEK polymer was calculated using where ∆*E* is the interaction energy at the interface between the CFs and the PEEK polymer, *Etotal* represents to the total potential energy of the composite, and *ECF* and *Epolymer* are the potential energy of the PI-sized CFs and PEEK polymer, respectively. The interfacial shear strength (ISS, *τ*) between the PI-sized CFs and the PEEK polymer was calculated using

where ∆ is the interaction energy at the interface between the CFs and the PEEK poly-

$$
\pi \left(MPa\right) = \frac{\Delta E}{WL^2} \tag{3}
$$

where is the ISS, is the width of the CF, represents the length of the CF, and ∆ corresponds to the difference in the interaction energy. where *τ* is the ISS, *W* is the width of the CF, *L* represents the length of the CF, and ∆*E* corresponds to the difference in the interaction energy.

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

#### *3.1. Synthesis and Characterization of PAS 3.1. Synthesis and Characterization of PAS*

The 1H NMR spectrum of PAS could be assigned by comparison between the 1H NMR spectra of HMP and PAA, as shown in Figure 4. The marked numbers and letters in the Figure 4 were noticed each molecular group and spectrum peak. The peaks at 2.56 ppm in the 1H NMR spectrum of HMP were related to the 3 protons of the CH3 groups. The protons in PAA could be confirmed from the peaks at 6–9 ppm in the 1H NMR spectrum of PAA. Those peaks were also found in the 1H NMR spectrum of PAS. Moreover, the 6 peaks at 1.5–4 ppm in the 1H NMR spectrum of PAS that appeared after the reaction of PAA with HMP were related to a proton on the piperidine ring of HMP. The proton of the CH3 group in HMP was de-shielded by a proton of the COOH group in PAA. The 1H NMR spectrum of PAS could be assigned by comparison between the 1H NMR spectra of HMP and PAA, as shown in Figure 4. The marked numbers and letters in the Figure 4 were noticed each molecular group and spectrum peak. The peaks at 2.56 ppm in the 1H NMR spectrum of HMP were related to the 3 protons of the CH3 groups. The protons in PAA could be confirmed from the peaks at 6–9 ppm in the 1H NMR spectrumof PAA. Those peaks were also found in the 1H NMR spectrum of PAS. Moreover, the 6 peaks at 1.5–4 ppm in the 1H NMR spectrum of PAS that appeared after the reaction of PAA with HMP were related to a proton on the piperidine ring of HMP. The proton of the CH3 group in HMP was de−shielded by a proton of the COOH group in PAA.

**Figure 4.** 1H NMR spectra of HMP, PAA, and PAS. **Figure 4.** 1H NMR spectra of HMP, PAA, and PAS.

Figure 5 shows the FT−IR spectra of PI, PAA, and PAS. The FT−IR spectrum of PAA shows N−H stretching at 3252 cm-1 , C=O stretching at 1620 cm-1 , N−H stretching at 1503 cm-1 , and C−O stretching at 1235 cm-1 . The FT−IR spectrum of PAS shows new bands at 3380 cm-1 and 2920 cm-1 , corresponding to the O−H and −CH3 bonds of the HMP groups in PAS. The FT−IR spectrum of PAS, which can also be observed in the FT−IR spectrum of PAA, indicates that PAS was completely synthesized, as shown in Figure 5. The FT−IR spectra of PAA, PAS, and PI were used to verify the imidization of PAS to obtain cured PI. This was conducted via the observation of the disappearance of the peak at 1620 cm-1 and the growth of a peak at 1720 cm-1 . In Figure 5, the photographic image represents the aqueous PAS solution in water with an increase in concentration according to our previous study [15]. The synthesized PAS powder was readily available to the water−dispersible sizing agent without organic solvents. Figure 5 shows the FT-IR spectra of PI, PAA, and PAS. The FT-IR spectrum of PAA shows N-H stretching at 3252 cm−1 , C=O stretching at 1620 cm−1 , N-H stretching at 1503 cm−1 , and C-O stretching at 1235 cm−1 . The FT-IR spectrum of PAS shows new bands at 3380 cm−1 and 2920 cm−1 , corresponding to the O-H and -CH3 bonds of the HMP groups in PAS. The FT-IR spectrum of PAS, which can also be observed in the FT-IR spectrum of PAA, indicates that PAS was completely synthesized, as shown in Figure 5. The FT-IR spectra of PAA, PAS, and PI were used to verify the imidization of PAS to obtain cured PI. This was conducted via the observation of the disappearance of the peak at 1620 cm−1 and the growth of a peak at 1720 cm−1 . In Figure 5, the photographic image represents the aqueous PAS solution in water with an increase in concentration according to our previous study [15]. The synthesized PAS powder was readily available to the water-dispersible sizing agent without organic solvents.

**Figure 5.** FT−IR spectra of PI, PAA, and PAS. **Figure 5.** FT-IR spectra of PI, PAA, and PAS.

#### *3.2. Surface Morphology Analysis of the PI−Sized CFs 3.2. Surface Morphology Analysis of the PI-Sized CFs*

The surface structures of the CFs were characterized using FT−IR, AFM, and FE−SEM observations. The FT−IR spectra of the C−CFs, de−CFs, and PI−sized CFs with contents of 0.5, 1.0, and 1.5 wt% were combined in Figure 6. The spectrum of the de−CFs was relatively smoother than that of the C−CFs due to the removal of epoxy groups and any other residues. A sharp carboxylic acid absorption peak (C=O stretching) in all the PI−sized CF samples was observed near 1700 cm-1 . The addition of PAS molecules onto the CF surfaces reveals N−H stretching. This was assigned to 1320 cm-1 , which was not displayed in the C−CFs. It could be determined that the carbon fiber surfaces were successfully modified with the cured PI from the newly synthesized water−dispersible PAS, which led to a more effective increase in the carbon fiber surface polarity and interfacial adhesion between the matrix and fibers. The 2D and 3D surface topographies of the de−CFs, C−CFs, and PI−sized CFs with different contents of PI sizing agents were compared, as shown in Figure 7. Increased surface roughness was observed on the PI−sized CF surfaces compared with those of the The surface structures of the CFs were characterized using FT-IR, AFM, and FE-SEM observations. The FT-IR spectra of the C-CFs, de-CFs, and PI-sized CFs with contents of 0.5, 1.0, and 1.5 wt% were combined in Figure 6. The spectrum of the de-CFs was relatively smoother than that of the C-CFs due to the removal of epoxy groups and any other residues. A sharp carboxylic acid absorption peak (C=O stretching) in all the PI-sized CF samples was observed near 1700 cm−1 . The addition of PAS molecules onto the CF surfaces reveals N-H stretching. This was assigned to 1320 cm−1 , which was not displayed in the C-CFs. It could be determined that the carbon fiber surfaces were successfully modified with the cured PI from the newly synthesized water-dispersible PAS, which led to a more effective increase in the carbon fiber surface polarity and interfacial adhesion between the matrix and fibers. *Polymers* **2023**, *15*, 1646 8 of 17

**Figure 6.** FT−IR spectra of C−CFs and de−CFs as well as 0.5, 1.0, and 1.5 wt% PI−sized CFs. **Figure 6.** FT-IR spectra of C-CFs and de-CFs as well as 0.5, 1.0, and 1.5 wt% PI-sized CFs.

**Figure 7.** AFM images and height profiles for (**a**) C−CFs and (**b**) de−CFs, as well as (**c**) PI 0.5 wt%,

(**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% of PI−sized CFs.

The 2D and 3D surface topographies of the de-CFs, C-CFs, and PI-sized CFs with different contents of PI sizing agents were compared, as shown in Figure 7. Increased surface roughness was observed on the PI-sized CF surfaces compared with those of the de-CFs. The average roughness (Ra) of the PI-sized CFs with contents of 0.5, 1.0, and 1.5 wt% (40.9, 62.1, and 44.0 nm, respectively) was greater than that of the de-CFs (32.5 nm). This increased roughness led to the enhancement of mechanical interlocking between the surface of the PI-sized CFs and the PEEK polymer [44]. However, in the case of the PI 1.5 wt% PI-sized CFs, the variation in height became relatively high (from −150 nm to 170 nm) due to an excessive amount of coating, as shown in Figure 7e [45,46]. This agglomerated non-uniform sizing layer between the carbon fibers and the matrix can lead to a decrease in interfacial adhesion. **Figure 6.** FT−IR spectra of C−CFs and de−CFs as well as 0.5, 1.0, and 1.5 wt% PI−sized CFs.

*Polymers* **2023**, *15*, 1646 8 of 17

**Figure 7.** AFM images and height profiles for (**a**) C−CFs and (**b**) de−CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% of PI−sized CFs. **Figure 7.** AFM images and height profiles for (**a**) C-CFs and (**b**) de-CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% of PI-sized CFs.

Observations using SEM-EDX were performed to analyze the compositional elements of reactive PI groups on the carbon fiber surfaces. Figure 8 shows the PI line profile of the EDX spectra across the carbon fiber surface along the scanline. The graph demonstrates the intensity of elemental carbon (C), nitrogen (N), and oxygen (O) with respect to distance along the fiber surface. The elements N and O had relatively uniform low concentrations over all of the carbon fiber surfaces. However, the C line profile shows an obvious decrease according to the increase in the content of the PI sizing agent, as shown in Figure 8a,e. This result was ascribed to a PI sizing layer containing uniformly attached N or O compared

with the C-CFs or de-CFs. Increasing the amount of N or O greatly affects the interfacial adhesion between CFs and a PEEK polymer due to solubility with higher ionization [47]. or O compared with the C−CFs or de−CFs. Increasing the amount of N or O greatly affects the interfacial adhesion between CFs and a PEEK polymer due to solubility with higher ionization [47].

Observations using SEM−EDX were performed to analyze the compositional elements of reactive PI groups on the carbon fiber surfaces. Figure 8 shows the PI line profile of the EDX spectra across the carbon fiber surface along the scanline. The graph demonstrates the intensity of elemental carbon (C), nitrogen (N), and oxygen (O) with respect to distance along the fiber surface. The elements N and O had relatively uniform low concentrations over all of the carbon fiber surfaces. However, the C line profile shows an obvious decrease according to the increase in the content of the PI sizing agent, as shown in Figure 8a,e. This result was ascribed to a PI sizing layer containing uniformly attached N

*Polymers* **2023**, *15*, 1646 9 of 17

**Figure 8.** FE−SEM images and SEM−EDX line spectra of elemental carbon (green line), nitrogen (blue line), and oxygen (red line) in (**a**) C−CFs and (**b**) de−CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% PI−sized CFs. **Figure 8.** FE-SEM images and SEM-EDX line spectra of elemental carbon (green line), nitrogen (blue line), and oxygen (red line) in (**a**) C-CFs and (**b**) de-CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% PI-sized CFs.

#### *3.3. Thermo−Mechanical Properties of the PI−Sized CF/PEEK Composites 3.3. Thermo-Mechanical Properties of the PI-Sized CF/PEEK Composites*

The ILSS results were analyzed to investigate the characteristics of the interface between the CFs and the PEEK polymer, with or without the PI sizing agent, as shown in Figure 9. The short−beam strength for the PI 0.5−, 1.0−, and 1.5−sized CF−reinforced PEEK composites was 76.9, 73.8, and 71.4 MPa, respectively. The ILSSs of all PI−sized CFs and de−CF/PEEK composites were higher (up to 13.2%) than those of the C−CF/PEEK composites with the commercial sizing agents. The structure of the aromatic PI synthesized from the PAA salt of the PMDA/ODA type shows a charge−transfer interaction consisting of electron donors (including nitrogen groups) and electron acceptors (including carbonyl groups). This charge−transfer complex was attributed to the intramolecular forces between PI molecules, as well as to the intermolecular forces between PI molecules and PEEK polymer chains. These properties were supported by the strong molecular attraction due to the formation of robust polymer chains. With all these reactive groups, it was possible for them to attach layer by layer. Therefore, polyimides may not only form well−stacked sizing layers on the surfaces of carbon fibers but can also interact well with the p−phenylene groups of PEEK polymer chains [31]. This leads to an enhanced potential for the PI sizing agent to cause improved interfacial adhesion between the CFs and the PEEK polymer. The fiber surfaces of the de−CFs could not be protected from abrasion due to the lack of the sizing agent. In addition, the de−CFs were difficult to handle due to low cohesion. Therefore, the de−CFs are limited in fiber processing and production and composite material fabrication. Therefore, it is necessary for the treatment with a sizing agent The ILSS results were analyzed to investigate the characteristics of the interface between the CFs and the PEEK polymer, with or without the PI sizing agent, as shown in Figure 9. The short-beam strength for the PI 0.5-, 1.0-, and 1.5-sized CF-reinforced PEEK composites was 76.9, 73.8, and 71.4 MPa, respectively. The ILSSs of all PI-sized CFs and de-CF/PEEK composites were higher (up to 13.2%) than those of the C-CF/PEEK composites with the commercial sizing agents. The structure of the aromatic PI synthesized from the PAA salt of the PMDA/ODA type shows a charge-transfer interaction consisting of electron donors (including nitrogen groups) and electron acceptors (including carbonyl groups). This charge-transfer complex was attributed to the intramolecular forces between PI molecules, as well as to the intermolecular forces between PI molecules and PEEK polymer chains. These properties were supported by the strong molecular attraction due to the formation of robust polymer chains. With all these reactive groups, it was possible for them to attach layer by layer. Therefore, polyimides may not only form well-stacked sizing layers on the surfaces of carbon fibers but can also interact well with the p-phenylene groups of PEEK polymer chains [31]. This leads to an enhanced potential for the PI sizing agent to cause improved interfacial adhesion between the CFs and the PEEK polymer. The fiber surfaces of the de-CFs could not be protected from abrasion due to the lack of the sizing agent. In addition, the de-CFs were difficult to handle due to low cohesion. Therefore, the de-CFs are limited in fiber processing and production and composite material fabrication. Therefore, it is necessary for the treatment with a sizing agent to be appropriate to preserve the carbon fibers [48,49]. Figure 9 shows that the 0.5 wt% content of the PI sizing agent improved the short-beam strength of the CF/PEEK composites by approximately 13%. The commercial sizing agent is generally known to achieve approximately 1% coating on the surfaces of carbon fibers. This result demonstrates that composites reinforced with the newly developed water-dispersible PI have competitiveness for their efficient application to high-performance structures.

to be appropriate to preserve the carbon fibers [48,49]. Figure 9 shows that the 0.5 wt% content of the PI sizing agent improved the short−beam strength of the CF/PEEK composites by approximately 13%. The commercial sizing agent is generally known to achieve approximately 1% coating on the surfaces of carbon fibers. This result demonstrates that composites reinforced with the newly developed water−dispersible PI have competitive-

ness for their efficient application to high−performance structures.

**Figure 9.** Interlaminar shear properties of composites reinforced with C−CFs and de−CFs, as well as PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% PI−sized CFs. (**a**) stress–strain curves; (**b**) interlaminar shear strengths and enhancements. **Figure 9.** Interlaminar shear properties of composites reinforced with C-CFs and de-CFs, as well as PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% PI-sized CFs. (**a**) stress–strain curves; (**b**) interlaminar shear strengths and enhancements.

The DMA results on the viscoelastic properties of the composites reinforced with the C−CFs, de−CFs, and PI−sized CFs are shown in Figure 10. These results indicate that the complex modulus between the polymers can be affected by a wide variety of chain motions over a range of temperatures. Figure 10a shows a comparative increase in stored energy for all of the composites containing PI−sized CFs compared with the composites containing commercial epoxy−sized CFs. The storage moduli of the composites with 0.5, 1.0, and 1.5 wt% PI−sized CFs were 23,076, 24,440, and 23,724 MPa, respectively. In the case of the 1.0 wt% PI−sized CFs, the rate of increase in the storage modulus was approximately 11.6% compared with that of the C−CF/PEEK composites. As mentioned above, this is because the interfacial adhesion between the carbon fibers and the PEEK polymer was improved due to enhanced intermolecular attraction by adding the PI sizing agent. The effect of the PI sizing agent was demonstrated while ensuring excellent thermal stability at a high temperature. In addition, the increase in the loss modulus, which indicates energy dissipated as heat from the composites, approaches that of its viscous property. The composite reinforced with 1.0 wt% PI−sized CFs is distinctively represented in Figure 10b. This phenomenon means that the intermolecular force associated with the interfacial interaction between the PI sizing agent and the PEEK polymer was even greater than the intramolecular force between the CFs and the PI sizing agent of the PI−sized CF/PEEK composites. This result indicates that the surface modification of the carbon fibers could be determined by the interfacial failure attributed to the physical properties of the composites. In addition, Figure 10c shows that the glass transition temperatures of the PI−sized CF/PEEK composites were increased in comparison with the C−CF/PEEK composites. Among the content levels of the PI sizing agent, the appropriate amount on the surface of a CF is 0.5 wt% for the storage modulus as well as the damping factor (tan δ). This means that the thermal resistance to molecular mobility in the composites with the PI sizing agent was enhanced in comparison with the C−CFs or de−CF/PEEK composites [34]. Therefore, the effective incorporation of PI on the CFs was important to improve the thermo−mechanical properties of the PEEK composites reinforced with CFs. As shown in Figure 10c, the measured width from the tan δ curve for the composite reinforced C−CFs is remarkably wide, which indicates a greater distribution of the segment lengths related The DMA results on the viscoelastic properties of the composites reinforced with the C-CFs, de-CFs, and PI-sized CFs are shown in Figure 10. These results indicate that the complex modulus between the polymers can be affected by a wide variety of chain motions over a range of temperatures. Figure 10a shows a comparative increase in stored energy for all of the composites containing PI-sized CFs compared with the composites containing commercial epoxy-sized CFs. The storage moduli of the composites with 0.5, 1.0, and 1.5 wt% PI-sized CFs were 23,076, 24,440, and 23,724 MPa, respectively. In the case of the 1.0 wt% PI-sized CFs, the rate of increase in the storage modulus was approximately 11.6% compared with that of the C-CF/PEEK composites. As mentioned above, this is because the interfacial adhesion between the carbon fibers and the PEEK polymer was improved due to enhanced intermolecular attraction by adding the PI sizing agent. The effect of the PI sizing agent was demonstrated while ensuring excellent thermal stability at a high temperature. In addition, the increase in the loss modulus, which indicates energy dissipated as heat from the composites, approaches that of its viscous property. The composite reinforced with 1.0 wt% PI-sized CFs is distinctively represented in Figure 10b. This phenomenon means that the intermolecular force associated with the interfacial interaction between the PI sizing agent and the PEEK polymer was even greater than the intramolecular force between the CFs and the PI sizing agent of the PI-Sized CF/PEEK composites. This result indicates that the surface modification of the carbon fibers could be determined by the interfacial failure attributed to the physical properties of the composites. In addition, Figure 10c shows that the glass transition temperatures of the PI-Sized CF/PEEK composites were increased in comparison with the C-CF/PEEK composites. Among the content levels of the PI sizing agent, the appropriate amount on the surface of a CF is 0.5 wt% for the storage modulus as well as the damping factor (tan δ). This means that the thermal resistance to molecular mobility in the composites with the PI sizing agent was enhanced in comparison with the C-CFs or de-CF/PEEK composites [34]. Therefore, the effective incorporation of PI on the CFs was important to improve the thermo-mechanical properties of the PEEK composites reinforced with CFs. As shown in Figure 10c, the measured width from the tan δ curve for the composite reinforced C-CFs is remarkably wide, which indicates a greater distribution of the segment lengths related to molecular mobility than in the PI-Sized CF/PEEK composites. Moreover, the intensity of the tan δ curve for the composite-reinforced C-CFs presents a greater increase in the mobility of the relaxing segments of the polymer compared with that of the composite-reinforced PI-sized CFs. This refers to the weak interaction bonding resulting from poor inter–intramolecular forces between the C-CFs and the PEEK polymer with increasing temperature. These results suggest that the interfacial interaction between CFs and PEEK can be affected by the PI sizing agent on the surfaces of the carbon fibers. This effect was attributed to the significantly enhanced interfacial adhesion of the composites. The effect of the sizing agent with high heat resistance was the remarkable thermal stability

of the CF/PEEK composites. In addition, the thermal stability of the PI-sized CFs with high heat resistance was compared with that of the C-CFs, as shown in Figure 11. The residual weights of the composites with C-CFs, de-CFs, PI0.5, PI1.0, and PI1.5 were 98.63, 99.55, 99.80, 99.59, and 99.67%, respectively. In the case of the PI-sized CFs with contents of 0.5, 1.0, and 1.5 wt%, there was no weight loss over the temperature range of 25 ◦C to 800 ◦C. The difference in the weight loss ratio between the PI-sized CFs and C-CFs was definitely observed depending on the absence of the polyimide sizing agent. Therefore, it was maintainable to protect the surfaces of the carbon fibers with the residual amount of the polyimide sizing agent after decomposition at a high temperature. CF/PEEK composites. In addition, the thermal stability of the PI−sized CFs with high heat resistance was compared with that of the C−CFs, as shown in Figure 11. The residual weights of the composites with C−CFs, de−CFs, PI0.5, PI1.0, and PI1.5 were 98.63, 99.55, 99.80, 99.59, and 99.67%, respectively. In the case of the PI−sized CFs with contents of 0.5, 1.0, and 1.5 wt%, there was no weight loss over the temperature range of 25 °C to 800 °C. The difference in the weight loss ratio between the PI−sized CFs and C−CFs was definitely observed depending on the absence of the polyimide sizing agent. Therefore, it was maintainable to protect the surfaces of the carbon fibers with the residual amount of the polyimide sizing agent after decomposition at a high temperature.

to molecular mobility than in the PI−sized CF/PEEK composites. Moreover, the intensity of the tan δ curve for the composite−reinforced C−CFs presents a greater increase in the mobility of the relaxing segments of the polymer compared with that of the composite−reinforced PI−sized CFs. This refers to the weak interaction bonding resulting from poor inter–intramolecular forces between the C−CFs and the PEEK polymer with increasing temperature. These results suggest that the interfacial interaction between CFs and PEEK can be affected by the PI sizing agent on the surfaces of the carbon fibers. This effect was attributed to the significantly enhanced interfacial adhesion of the composites. The effect of the sizing agent with high heat resistance was the remarkable thermal stability of the

*Polymers* **2023**, *15*, 1646 11 of 17

**Figure 10.** (**a**) Storage moduli, (**b**) loss moduli, and (**c**) tan δ curves for the composites reinforced with C−CFs and de−CFs, as well as PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% of PI−sized CFs. **Figure 10.** (**a**) Storage moduli, (**b**) loss moduli, and (**c**) tan δ curves for the composites reinforced with C-CFs and de-CFs, as well as PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% of PI-sized CFs.

**Figure 11.** TGA curves of C−CFs, de−CFs, PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% sized CFs. **Figure 11.** TGA curves of C-CFs, de-CFs, PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% sized CFs. **Figure 11.** TGA curves of C−CFs, de−CFs, PI 0.5 wt%, PI 1.0 wt%, and PI 1.5 wt% sized CFs.

#### *3.4. Fracture Characteristics of the CF/PEEK Composites 3.4. Fracture Characteristics of the CF/PEEK Composites 3.4. Fracture Characteristics of the CF/PEEK Composites*

Figure 12 shows SEM images of the fracture surfaces on the composites reinforced with carbon fibers after the ILSS test. The fracture surfaces of the composites reinforced with C−CFs and de−CFs are relatively rough compared with those of the composites reinforced with PI−sized CFs at all content levels. Damaged fibers in the fill direction were rumpled, and pull−out of fibers were observed in the warp direction (Figure 12a,b). In addition, while delamination between the C−CFs or de−CFs and PEEK matrix was observed, the fracture surfaces of the PI 0.5, 1.0, and 1.5 wt% PEEK composites are rather smooth and tightly embedded between the CFs and the PEEK matrix. This phenomenon confirms that the interfacial adhesion was enhanced by adding the PI sizing agent to the carbon fibers [50]. The PI sizing agent was protected due to its high heat resistance during high−temperature manufacturing processes, which allowed it to withstand the effective stress transfer of the carbon fiber. The PI−sized CFs played an important role in improving the interfacial performance of these PEEK composites. Figure 12 shows SEM images of the fracture surfaces on the composites reinforced with carbon fibers after the ILSS test. The fracture surfaces of the composites reinforced with C-CFs and de-CFs are relatively rough compared with those of the composites reinforced with PI-sized CFs at all content levels. Damaged fibers in the fill direction were rumpled, and pull-out of fibers were observed in the warp direction (Figure 12a,b). In addition, while delamination between the C-CFs or de-CFs and PEEK matrix was observed, the fracture surfaces of the PI 0.5, 1.0, and 1.5 wt% PEEK composites are rather smooth and tightly embedded between the CFs and the PEEK matrix. This phenomenon confirms that the interfacial adhesion was enhanced by adding the PI sizing agent to the carbon fibers [50]. The PI sizing agent was protected due to its high heat resistance during high-temperature manufacturing processes, which allowed it to withstand the effective stress transfer of the carbon fiber. The PI-sized CFs played an important role in improving the interfacial performance of these PEEK composites. Figure 12 shows SEM images of the fracture surfaces on the composites reinforced with carbon fibers after the ILSS test. The fracture surfaces of the composites reinforced with C−CFs and de−CFs are relatively rough compared with those of the composites reinforced with PI−sized CFs at all content levels. Damaged fibers in the fill direction were rumpled, and pull−out of fibers were observed in the warp direction (Figure 12a,b). In addition, while delamination between the C−CFs or de−CFs and PEEK matrix was observed, the fracture surfaces of the PI 0.5, 1.0, and 1.5 wt% PEEK composites are rather smooth and tightly embedded between the CFs and the PEEK matrix. This phenomenon confirms that the interfacial adhesion was enhanced by adding the PI sizing agent to the carbon fibers [50]. The PI sizing agent was protected due to its high heat resistance during high−temperature manufacturing processes, which allowed it to withstand the effective stress transfer of the carbon fiber. The PI−sized CFs played an important role in improving the interfacial performance of these PEEK composites.

**Figure 12.** FE−SEM images of fracture surfaces on composites reinforced with (**a**) C−CFs and (**b**) de−CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% of the PI−sized CFs. **Figure 12.** FE−SEM images of fracture surfaces on composites reinforced with (**a**) C−CFs and (**b**) de−CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% of the PI−sized CFs. **Figure 12.** FE-SEM images of fracture surfaces on composites reinforced with (**a**) C-CFs and (**b**) de-CFs, as well as (**c**) PI 0.5 wt%, (**d**) PI 1.0 wt%, and (**e**) PI 1.5 wt% of the PI-sized CFs.

#### *3.5. Molecular Dynamics of the CF/PEEK Composite 3.5. Molecular Dynamics of the CF/PEEK Composite*

MSD analysis can derive an estimation of the parameters of molecular movements, such as the diffusion coefficient. When the model systems reached the phase transition region, the diffusion level increased sharply, which can determine the glass transition temperature. The temperatures of the phase transition regions of the composite models were increased, and are marked with a red arrow in Figure 13. The temperatures of the phase transition regions of the composite models with de-CFs, PI0.3, and PI1.0 were 400–450 K, 425–475 K, and 450–475 K, respectively. The de-CFs (Figure 13a) show a relatively low temperature range compared with the PI-sized composites. This result shows that the PI sizing agent increased the glass transition temperature. MSD analysis can derive an estimation of the parameters of molecular movements, such as the diffusion coefficient. When the model systems reached the phase transition region, the diffusion level increased sharply, which can determine the glass transition temperature. The temperatures of the phase transition regions of the composite models were increased, and are marked with a red arrow in Figure 13. The temperatures of the phase transition regions of the composite models with de−CFs, PI0.3, and PI1.0 were 400–450 K, 425–475 K, and 450–475 K, respectively. The de−CFs (Figure 13a) show a relatively low temperature range compared with the PI−sized composites. This result shows that the PI sizing agent increased the glass transition temperature.

**Figure 13.** Mean square displacement (MSD) results: (**a**) de−CFs, (**b**) PI0.3, and (**c**) PI1.0. **Figure 13.** Mean square displacement (MSD) results: (**a**) de-CFs, (**b**) PI0.3, and (**c**) PI1.0.

The interfacial characteristics between the de−CFs and the PEEK polymer were calculated using MD simulations. Table 1 shows the interaction energy (∆) between the CFs and the PEEK polymer. The negative interfacial energy means that the energy of adsorption of the PEEK molecules at the interface was strong [51,52]. Therefore, the high negative value indicates strong adhesion at the interface between the CFs and the PEEK polymer. The interaction energy of the de−CF/PEEK composites was -1788.79 kcal/mol. Those of the PI−sized CF/PEEK composites with contents of PI 0.3, 1.0, 2.0, and 5.0 wt% were -1877.73, -1869.48, -1815.94, and -1767.29 kcal/mol, respectively. These results show that the interaction energies of the PI−sized CF/PEEK composites with contents of PI 0.3, 1.0, and 2.0 wt% generally had a higher interaction energy than that of the de−CF/PEEK composite. In addition, the PI sizing agents also affected the interfacial shear strength (ISS), as shown in Figure 14. The interfacial shear strengths of the de−CFs, PI0.3, PI1.0, PI2.0, and PI5.0 were calculated to be 103.15, 107.17, 107.31, 103.84, and 101.24 MPa, respectively, obtained from Equation (3). The simulated interfacial shear strengths of the PI−sized CF/PEEK composites, except for the PI5.0 model, were higher than that of the de−CF/PEEK composite. This result proves that the PI sizing agent affected the interfacial properties between the CFs and the PEEK polymer. In addition, the contents of 0.3 wt% and 1.0 wt% of the PI sizing agent on the CFs were the most effective for improving the interaction energy and interfacial shear strength. The results obtained from the MD simulations show similar trends to the experimental data. The interfacial characteristics between the de-CFs and the PEEK polymer were calculated using MD simulations. Table 1 shows the interaction energy (∆*E*) between the CFs and the PEEK polymer. The negative interfacial energy means that the energy of adsorption of the PEEK molecules at the interface was strong [51,52]. Therefore, the high negative value indicates strong adhesion at the interface between the CFs and the PEEK polymer. The interaction energy of the de-CF/PEEK composites was −1788.79 kcal/mol. Those of the PI-Sized CF/PEEK composites with contents of PI 0.3, 1.0, 2.0, and 5.0 wt% were −1877.73, −1869.48, −1815.94, and −1767.29 kcal/mol, respectively. These results show that the interaction energies of the PI-Sized CF/PEEK composites with contents of PI 0.3, 1.0, and 2.0 wt% generally had a higher interaction energy than that of the de-CF/PEEK composite. In addition, the PI sizing agents also affected the interfacial shear strength (ISS), as shown in Figure 14. The interfacial shear strengths of the de-CFs, PI0.3, PI1.0, PI2.0, and PI5.0 were calculated to be 103.15, 107.17, 107.31, 103.84, and 101.24 MPa, respectively, obtained from Equation (3). The simulated interfacial shear strengths of the PI-Sized CF/PEEK composites, except for the PI5.0 model, were higher than that of the de-CF/PEEK composite. This result proves that the PI sizing agent affected the interfacial properties between the CFs and the PEEK polymer. In addition, the contents of 0.3 wt% and 1.0 wt% of the PI sizing agent on the CFs were the most effective for improving the interaction energy and interfacial shear strength. The results obtained from the MD simulations show similar trends to the experimental data.

**Table 1.** Interaction energy of CF/PEEK composites using MD simulation. **Table 1.** Interaction energy of CF/PEEK composites using MD simulation.

**Figure 14.** Interfacial shear strength of PEEK composites reinforced with de−CFs as well as PI 0.3 wt%, PI 1.0 wt%, PI 2.0 wt%, and PI 5.0 wt% sized CFs using MD simulation. **Figure 14.** Interfacial shear strength of PEEK composites reinforced with de-CFs as well as PI 0.3 wt%, PI 1.0 wt%, PI 2.0 wt%, and PI 5.0 wt% sized CFs using MD simulation.

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

sites in the future.

In this study, the effects of water−dispersible PI−sized CFs on the thermo−mechanical properties of PEEK composites were investigated. The PMDA/ODA used for the PAA solution was successfully synthesized and thermally imidized to apply it as a sizing agent for the carbon fibers. The surfaces of the carbon fibers were modified by adding a PI sizing layer that led to improved interfacial adhesion between the CFs and PEEK matrix. The short−beam strengths and storage moduli of the PEEK composites reinforced with PI−sized CFs increased more than those of the composites reinforced with de−sized and commercial CFs. The surface roughness of the PI−sized CFs as indicated in the 2D and 3D topographies was greater than those of the C−CFs and de−CFs. This result shows that the CF surfaces were coated with the PI sizing agent. The fracture surfaces of the composites reinforced with PI−sized CFs exhibited greater smoothness and were more tightly embedded between the CFs and PEEK matrix than those of the composites reinforced with C−CFs and de−CFs. This was ascribed to enhanced mechanical interlocking between the PI−sized CFs and the PEEK polymer of the composites. These results suggest that the interfacial interaction between CFs and PEEK can be affected by a PI sizing agent on the surfaces of carbon fibers. This effect was attributed to the significantly enhanced interfacial adhesion of the composites even at a high temperature. MD simulations were performed to examine the effects of the PI sizing agents on the interfacial properties between the CFs and the PEEK polymer. The interfacial properties between the CFs and the PEEK polymer were most improved when the PI sizing agent with 0.3 wt% was applied to the surface of the carbon fibers. As a result, the newly developed PI sizing agent with high heat resistance can be used to improve the interfacial bonding properties and thermo−mechanical properties of CF/PEEK composites. The present study focused on improving the interfacial bonding properties and thermo−mechanical properties of CF/PEEK composites. PI−sized carbon nanomaterials with high−temperature resistance and high dispersion can In this study, the effects of water-dispersible PI-sized CFs on the thermo-mechanical properties of PEEK composites were investigated. The PMDA/ODA used for the PAA solution was successfully synthesized and thermally imidized to apply it as a sizing agent for the carbon fibers. The surfaces of the carbon fibers were modified by adding a PI sizing layer that led to improved interfacial adhesion between the CFs and PEEK matrix. The short-beam strengths and storage moduli of the PEEK composites reinforced with PI-sized CFs increased more than those of the composites reinforced with de-sized and commercial CFs. The surface roughness of the PI-sized CFs as indicated in the 2D and 3D topographies was greater than those of the C-CFs and de-CFs. This result shows that the CF surfaces were coated with the PI sizing agent. The fracture surfaces of the composites reinforced with PI-sized CFs exhibited greater smoothness and were more tightly embedded between the CFs and PEEK matrix than those of the composites reinforced with C-CFs and de-CFs. This was ascribed to enhanced mechanical interlocking between the PI-sized CFs and the PEEK polymer of the composites. These results suggest that the interfacial interaction between CFs and PEEK can be affected by a PI sizing agent on the surfaces of carbon fibers. This effect was attributed to the significantly enhanced interfacial adhesion of the composites even at a high temperature. MD simulations were performed to examine the effects of the PI sizing agents on the interfacial properties between the CFs and the PEEK polymer. The interfacial properties between the CFs and the PEEK polymer were most improved when the PI sizing agent with 0.3 wt% was applied to the surface of the carbon fibers. As a result, the newly developed PI sizing agent with high heat resistance can be used to improve the interfacial bonding properties and thermo-mechanical properties of CF/PEEK composites. The present study focused on improving the interfacial bonding properties and thermo-mechanical properties of CF/PEEK composites. PI-sized carbon nanomaterials with high-temperature resistance and high dispersion can be investigated to improve thethermo-mechanical properties of polymer-based composites in the future.

be investigated to improve the thermo−mechanical properties of polymer−based compo-

**Author Contributions:** Conceptualization, H.J. and K.J.B.; methodology, H.J. and J.-U.J.; software, K.J.B. and Y.O.; validation, N.-H.Y. and J.Y.; formal analysis, K.J.B. and H.J.; investigation, K.J.B.; data curation, H.J.; writing—original draft preparation, H.J. and K.J.B.; writing—review and editing, **Author Contributions:** Conceptualization, H.J. and K.J.B.; methodology, H.J. and J.-U.J.; software, K.J.B. and Y.O.; validation, N.-H.Y. and J.Y.; formal analysis, K.J.B. and H.J.; investigation, K.J.B.; data curation, H.J.; writing-original draft preparation, H.J. and K.J.B.; writing-review and editing, N.-H.Y. and J.Y.; supervision, J.Y.; project administration, J.Y.; funding acquisition, J.Y., H.J. and K.J.B. All authors contributed equally to this paper. All authors have read and agreed to the published version of the manuscript.

**Funding:** This study was supported by the Korea Institute of Science and Technology (KIST) Institutional Program. This study was also supported by the Technology Innovation Program 'Development of an ultra-lightweight 19 carbon composite wheel using integrated braid preform manufacturing technology with tow prepreg material' (no. 20021913) funded by the Ministry of Trade, Industry, and Energy (MOTIE, Korea).

**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:** This study was supported by the Korea Institute of Science and Technology (KIST) Institutional Program.

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

### **References**

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## *Article* **Effect of Short Fibers on Fracture Properties of Epoxy-Based Polymer Concrete Exposed to High Temperatures**

**Oussama Elalaoui 1,2**

**Abstract:** Recently, polymer concrete (PC) has been widely used in many civil engineering applications. PC shows superiority in major physical, mechanical, and fracture properties comparing to ordinary Portland cement concrete. Despite many suitable characteristics of thermosetting resins related to processing, the thermal resistance of polymer concrete composite is relatively low. This study aims to investigate the effect of incorporating short fibers on mechanical and fracture properties of PC under different ranges of high temperatures. Short carbon and polypropylene fibers were added randomly at a rate of 1 and 2% by the total weight of the PC composite. The exposure temperatures cycles were ranged between 23 to 250 ◦C. Various tests were conducted including flexure strength, elastic modulus, toughness, tensile crack opening, density, and porosity to evaluate the effect of addition of short fibers on fracture properties of PC. The results show that the inclusion of short fiber lead to an increase in the load carrying capacity of PC by an average of 24% and limits the crack propagation. On the other hand, the enhancement of fracture properties of based PC containing short fibers is vanished at high temperature (250 ◦C), but still more efficient than ordinary cement concrete. This work could lead to broader applications of polymer concrete exposed to high temperatures.

**Keywords:** polymer concrete; short fibers; fracture properties; thermal resistance; porosity

**1. Introduction**

Polymer concrete (PC), one of the three types of polymer concrete type, is formed by completely substituting Portland cement with polymer binders. The two other types are polymer-impregnated concrete (PIC), which consists of a hardened Portland cement concrete impregnated with a low-viscosity monomer polymerized in situ, and the second one is the polymer cement concrete (PCC) which is produced by incorporating a monomer or polymer in a Portland cement concrete mix and polymerized after placing concrete [1].

The polymer binder used in PC is acting through a polymerization process to bond the solid particles of concrete [2,3]. PC, whether reinforced or non-reinforced, shows superiority in major physical, mechanical, and fracture properties when compared to ordinary Portland cement. These properties include high mechanical strength, rapid hardening, and short curing time, which are very beneficial for precast element production or partitions with high toughness. [4]. In addition, PC is attributed with fast demolding, minimum cracking, higher mechanical strength, good bonding to old substrates, better abrasion resistance, chemical resistance, and less maintenance in comparison to Portland cement concrete [5,6]. This explains the quick development of using PC in many civil engineering applications [7].

Various types of resins have been used in the production of polymer concrete such as epoxy [8,9], unsaturated polyester [10,11], acrylics and vinyl-ester, poly acrylate, polypropylene, furan-based resin, and other polymers [12,13]. However, the formed composites incorporating these materials are brittle and the majority of their properties are very sensitive to long exposure at high temperatures [14–17]. This is because the polymers acting as

**Citation:** Elalaoui, O. Effect of Short Fibers on Fracture Properties of Epoxy-Based Polymer Concrete Exposed to High Temperatures. *Polymers* **2023**, *15*, 1078. https:// doi.org/10.3390/polym15051078

Academic Editors: S. D. Jacob Muthu and Victor Tcherdyntsev

Received: 29 December 2022 Revised: 25 January 2023 Accepted: 2 February 2023 Published: 21 February 2023

**Copyright:** © 2023 by the author. 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/).

binders in the PC are organic substances, which exhibits lower resistant to heat than the inorganic ones. Accordingly, PC is not recommended to be used in prolonged high thermal exposure to due to it generates the degradation of the resin, which possibly results in a loss of mechanical and fracture strength. These drawbacks are limiting the extensive use of PC as becoming not more efficient option [18,19].

Several studies have been undertaken to improve the PC performance throughout the insertion of flame retardant, coupling agent, and fibers, or by the replacement of the solid part by industrial by-products (i.e., industrial wastes, recycle aggregates, etc.) [2,4,20–23]. However, few studies have been performed to investigate the retaining of mechanical and fracture properties of PC after long exposure in extreme thermal conditions. Elalaoui et al. showed that the mechanical and physical PC containing epoxy was obviously influenced by thermal conditions but still more efficient than Portland cement concrete even after exposure to temperatures around 250 ◦C [13]. Apart from the mechanical and fracture properties, PC composite incorporating epoxy as main binder material are relatively more expensive than composite made with other thermosets such as polyesters [9,24]. Regardless their high cost ranging between three and fifteen times more than other thermosets, epoxy resins are preferred in special applications where high cost can be easily balanced by their overall superior mechanical and fatigue properties [7,11,25,26].

The inclusion of synthetic or natural fibers seems to be an effective option to overcome the PC shortcoming and increase the strength capacity, ductility, and toughness of PC [27,28]. G.Martínez et al. [27] investigated the effects of adding polypropylene fibers on compressive properties of polymer concrete. Their results revealed that compressive of PC reinforced with fibers was improved. Reis et al. studied the effect of reinforcing epoxy PC by chopped glass (inorganic) or carbon fibers (organic) with different proportions [11,29]. A similar approach was taken by Rokbi et al. [30] who tried to valorize the use of natural resources by using laminated vegetable fibers consisting of woven fabric jute in various orientations to lead up the reduction in the environmental impact and improve the mechanical properties of PC. In both studies, it was shown clearly that the fractures, elasticity modulus, compressive, and splitting tensile and flexural strength of PC reinforced with fibers were improved.

Many models have been used to predict the crack propagation for different materials [29] such as non-linear elastic fracture mechanics, linear elastic fracture mechanics (LEFM), fictitious crack model (FCM), effective crack model (ECM), cohesive crack model (CCM), size-effect model, and two-parameter model (TPM). It was reported that the carbon fiber reinforcement improves the toughness of the PC by 29%, while the glass fibers generate a smaller improvement of 13%. The tenacity of PC is 36% higher than that of Portland cement concrete which is between 0.74 and 1.53 [11,30]. The last finding indicates that the PC is more resistant to crack widening even without reinforcements.

The TPM developed by Jenq and Shah [31] and based on determining the change in potential energy when crack extents will be used in this study as it is considered as one of the fracture models to implement [11].

A review of recent literature showed a considerable number of studies on polymer PC. However, limited work has been conducted to investigate the mechanical and fracture properties of fiber reinforced polymer concrete, which undergoes extreme thermal exposure conditions. Hence, the main objective of this study is to evaluate the degradation and the retention of the residual mechanical characteristics and toughness of epoxy-based polymer concrete reinforced randomly with short carbon and polypropylene fibers after exposure to elevated temperatures. In this study, short fibers were added in a rate of 1 and 2% by the total weight of the PC composite. The exposure temperatures cycles used were ranged between 23 to 250 ◦C. Various tests were conducted including flexure strength, elastic modulus, toughness, tensile crack opening, density, and porosity to evaluate the effect of addition of short fibers on fracture properties of PC.

### **2. Materials, Sample Preparation and Methods**

### *2.1. Materials*

A thermosetting two-component product (hardener and resin) was selected to prepare the polymer concrete. Commercialized epoxy resin under the reference EPONAL 371 obtained from Bostik SA (France) was used in this study. The epoxy was cross-linked with chemical modified polyamines hardener. It was mixed immediately before use for common applications. Advantageously, the curing of the selected thermosetting resin can be implemented at a temperature at 10–35 ◦C, but preferably at room temperatures (20–25 ◦C) as prescribed. The service temperature range of this epoxy has not been mentioned in technical data but the glass transition temperature was 82.51 ◦C, as measured using a differential scanning calorimetry machine TA Instruments Q100 DSC (Waters LLC, New Castle, UK). The epoxy properties are presented in Table 1.

**Table 1.** Properties of used materials according to the manufacturer.

Two kinds of fibers (carbon and polypropylene) were chosen to be embedded into the PC's mixes (Table 1). A carbon fiber marketed as Aceca® ECO-H6 was supplied by Torayca Company, Japan. These fibers (having cylindrical shapes) are chopped into short lengths in the interest to help dispersion in viscous fluids. The second type of fiber, labeled under the name of DUOMIX FIRE® (M6), was provided by Bekaert (Belgium) as fine monofilaments polypropylene. These fibers have a tubular shape and come in the form of clusters which can be easily dispersed during mixing. The main disadvantage of these types of fibers was lowering the workability of the mix for the same volume of polymer.

Sand graded from 0 to 4 mm and a gravel graded from 4 to 10 mm were used as fine and coarse aggregates, respectively. The specific gravity of the sand was 2.47, whereas the specific gravity of gravel was 2.53 g/cm3 . Preliminary investigations based on the compressible packing model (CPM) was conducted to obtain the proportions of aggregates that guaranteed the maximum compactness of solid grains while minimizing the amount of polymer needed to wet fully the aggregates and fill gaps between. Portland cement CEM I 52.5 N having a specific density of 3.13 g/cm3 and a specific surface equal to 3590 cm2/g as evaluated via Blaine apparatus was also used in this study.

### *2.2. Sample Preparation*

In this study, four PC mixes with fibers reinforcements and 13% of polymer, by weight of PC, were prepared. The PC samples were prepared according to a process described in the work of Elalaoui et al. [13].

The content of polymer was kept constant for the all mixes. This content was selected as the optimal amount in the basis of previous studies [13,25,32–35]. The fibers were added at the rate of 1% and 2%, by the total weight of the polymer. The PC compositions are listed

in Table 2. The formed PC mixes are named as PC-Xp, where Xp% denotes the type of reinforcement (C for carbon and P for polypropylene) and p referring to the mass fraction of the introduced fibers in term of percentage. The PC was casted according to a well-defined process as described in a previous published study [13]. Ordinary Portland cement concrete (OCC) samples were also casted by mixing the same aggregates proportions for comparison purposes. The OCC mix constituent is shown in Table 2. at the rate of 1% and 2%, by the total weight of the polymer. The PC compositions are listed in Table 2. The formed PC mixes are named as PC-Xp, where Xp% denotes the type of reinforcement (C for carbon and P for polypropylene) and p referring to the mass fraction of the introduced fibers in term of percentage. The PC was casted according to a welldefined process as described in a previous published study [13]. Ordinary Portland cement concrete (OCC) samples were also casted by mixing the same aggregates propor-

tions for comparison purposes. The OCC mix constituent is shown in Table 2.

The content of polymer was kept constant for the all mixes. This content was selected as the optimal amount in the basis of previous studies [13,25,32–35]. The fibers were added

**Table 2.** PC and OCC mixes with the varying fiber's content.

*Polymers* **2023**, *15*, x 4 of 16

#### *2.3. Testing Method 2.3. Testing Method*

Several tests were conducted including compression and flexure strength, elastic modulus determined using the ultrasonic pulse velocity method (UPV), toughness, tensile crack opening, density, and porosity (using mercury intrusion porosimetry, MIP) to evaluate the effect of addition of short fibers on fracture properties of PC. The principle of MIP is to inject a no-wetting liquid such as mercury inside the pores of a small sample (10 cm2 as maximum) under vacuum conditions and under pressure ranging from 14 to 414 MPa. Several tests were conducted including compression and flexure strength, elastic modulus determined using the ultrasonic pulse velocity method (UPV), toughness, tensile crack opening, density, and porosity (using mercury intrusion porosimetry, MIP) to evaluate the effect of addition of short fibers on fracture properties of PC. The principle of MIP is to inject a no-wetting liquid such as mercury inside the pores of a small sample (10 cm2 as maximum) under vacuum conditions and under pressure ranging from 14 to 414 MPa.

A uniaxial compression tests were carried out using cylinders sample (50 mm × 100 mm). The effect of adding fibers on the flexure strength, toughness, and mechanical residual properties of PC were investigated by performing three-point bending test beam (50 mm × 50 mm × 305 mm). These tests were conducted in a displacement-control mode with a constant rate of 1.0 mm/min and 1.25 mm/min for bending and compression, respectively, as per the procedure stated in RILEM PCM8 1995 [36]. A uniaxial compression tests were carried out using cylinders sample (50 mm × 100 mm). The effect of adding fibers on the flexure strength, toughness, and mechanical residual properties of PC were investigated by performing three-point bending test beam (50 mm × 50 mm × 305 mm). These tests were conducted in a displacement-control mode with a constant rate of 1.0 mm/min and 1.25 mm/min for bending and compression, respectively, as per the procedure stated in RILEM PCM8 1995 [36].

The midpoint deflection of beams was evaluated by the mean of a linear variable differential transformer (LVTD) rested on a steel L-shape bracket attached to specimens at mid-height as shown in the figure below (Figure 1). The midpoint deflection of beams was evaluated by the mean of a linear variable differential transformer (LVTD) rested on a steel L-shape bracket attached to specimens at mid-height as shown in the figure below (Figure 1).

**Figure 1.** Details of bending test apparatus. **Figure 1.** Details of bending test apparatus.

Concrete toughness was evaluated throughout pre-notched beams using a 250 kN closed loop INSTRON machine. The crack mouth opening displacement (CMOD) was measured by virtue of clip-gauge attached to a knife-edges installed at the bottom of the beams (Figure 2). The mid-span deflection was recorded using an LVDT placed at the left side at the bottom of the specimens. The dimensions of sample and central U-shape notch properties are itemized in Table 3. measured by virtue of clip-gauge attached to a knife-edges installed at the bottom of the beams (Figure 2). The mid-span deflection was recorded using an LVDT placed at the left side at the bottom of the specimens. The dimensions of sample and central U-shape notch properties are itemized in Table 3.

Concrete toughness was evaluated throughout pre-notched beams using a 250 kN closed loop INSTRON machine. The crack mouth opening displacement (CMOD) was

**Figure 2.** Experimental setup for CMOD controlled three-point bending test. **Figure 2.** Experimental setup for CMOD controlled three-point bending test.

*Polymers* **2023**, *15*, x 5 of 16

For the OCC mix, beam samples of 70 mm × 70 mm × 280 mm for bending test and **Table 3.** Dimensions of sample and central notch properties.

**Rete of Loading**

**b × d × L (mm/min) S Notch ao Notch** *d* PC 50 × 70 × 480 0.05 400 14 3 max 0.200 OCC 80 × 150 × 750 0.05 640 48 3 max 0.312 For the OCC mix, beam samples of 70 mm × 70 mm × 280 mm for bending test and cylinder samples of 150 mm × 300 mm for compression sample were casted. The OCC beams samples were loaded at rate of 0.05 MPa/s, whereas the cylinder samples was loaded at 0.5 MPa/s.

**Span Length**

cylinder samples of 150 mm × 300 mm for compression sample were casted. The OCC

The fracture toughness values of PC in mode-I, KIC, and the fracture energy GF were calculated using a two fracture parameters model [37,38] which appears to give rather realistic prediction of concrete fracture behavior. An actual crack is replaced by an equivalent fictitious crack [39]. The fracture toughness values of PC in mode-I, KIC, and the fracture energy GF were calculated using a two fracture parameters model [37,38] which appears to give rather realistic prediction of concrete fracture behavior. An actual crack is replaced by an equivalent fictitious crack [39].

The effective crack length and the peak load, are required to determine KIc using linear elastic fracture mechanics, LEFM. This value of KIc was found to be independent of the specimen size [40]. The effective crack length and the peak load, are required to determine KIc using linear elastic fracture mechanics, LEFM. This value of KIc was found to be independent of the specimen size [40].

The stress intensity factor is given by means of Equation (1): The stress intensity factor is given by means of Equation (1):

$$\mathbf{K\_{lc}} = \sigma\_{\rm NC} \sqrt{\pi \mathbf{a\_c}} \,\mathrm{F\_1(\alpha)}\tag{1}$$

where σNC= 3F.S 2b.d 2 and α= ac+d0 d+d0 where σNC = 3F.S 2b.d2 and α = ac+d0 d+d0

**Beam Dimensions**

$$F\_1(\alpha) = \frac{1.83 - 1.85\alpha + 4.76\alpha^2 - 5.3\alpha^3 + 2.51\alpha^4}{(1 + 2\alpha)(1 - \alpha)^{\frac{3}{2}}} \tag{2}$$

**Length of the**

**Width of the**

0 *a*

The critical crack length is presented by Equation (3) The critical crack length is presented by Equation (3)

$$\mathbf{a}\_{\mathbf{c}} = \mathbf{a}\_{0} \frac{\mathbf{C}\_{\mathbf{u}} \, f\_{2}(\alpha\_{0})}{\mathbf{C}\_{\mathbf{i}} \, f\_{2}(\alpha)} \text{ and } \mathbf{a}\_{0} \frac{\mathbf{a}\_{0}}{\mathbf{d}} \tag{3}$$

2

The geometrical parameters S, b, d, a0 and d0 are described in Figure 3.

The geometrical parameters S, b, d, a0 and d0 are described in Figure 3.

*Polymers* **2023**, *15*, x 6 of 16

**Figure 3.** Determination of fracture parameters of concrete (**a**) geometrically characteristic values, (**b**) typical F-CMOD curve **Figure 3.** Determination of fracture parameters of concrete (**a**) geometrically characteristic values, (**b**) typical F-CMOD curve. (**b**) typical F-CMOD curve The fracture energy is calculated according to the following relation as postulated by

The geometrical parameters S, b, d, a0 and d0 are described in Figure 3.

The fracture energy is calculated according to the following relation as postulated by the following LEFM equation: The fracture energy is calculated according to the following relation as postulated by the following LEFM equation: the following LEFM equation:

$$\mathbf{K\_{lc}} = \sqrt{\mathbf{E.G\_F}} \tag{4}$$

$$\mathbf{\color{red}{\cdot\$}} \qquad \qquad \qquad \qquad \qquad \qquad \qquad \dots \qquad \qquad \dots \qquad \qquad \dots$$

(4)

The prepared specimens were heated at a rate of 0.5 °C/min until target temperature was reached according to the predefined protocol. The temperature was maintained for 3 h before being decreases with the same heating rate till samples reached the ambient temperature being kept for 24 h until the specimens were tested. To ensure that the setting temperature (from 100 °C to 250 °C) is compatible with specimen temperature, two k-type thermocouples mounted on data logger were attached to one side and in the middle of The prepared specimens were heated at a rate of 0.5 ◦C/min until target temperature was reached according to the predefined protocol. The temperature was maintained for 3 h before being decreases with the same heating rate till samples reached the ambient temperature being kept for 24 h until the specimens were tested. To ensure that the setting temperature (from 100 ◦C to 250 ◦C) is compatible with specimen temperature, two k-type thermocouples mounted on data logger were attached to one side and in the middle of samples for a continuous recording (Figure 4). The prepared specimens were heated at a rate of 0.5 °C/min until target temperature was reached according to the predefined protocol. The temperature was maintained for 3 h before being decreases with the same heating rate till samples reached the ambient temperature being kept for 24 h until the specimens were tested. To ensure that the setting temperature (from 100 °C to 250 °C) is compatible with specimen temperature, two k-type thermocouples mounted on data logger were attached to one side and in the middle of samples for a continuous recording (Figure 4).

**Figure 4.** Temperatures control through installation of thermocouples. **Figure 4.** Temperatures control through installation of thermocouples.

A set of four specimens was allocated for each test type. The tests were conducted on concretes allowed to cure for 7 days at prescribed room temperature for both unreinforced A set of four specimens was allocated for each test type. The tests were conducted on concretes allowed to cure for 7 days at prescribed room temperature for both unreinforced and reinforced PC and 28 days age for OCC. A set of four specimens was allocated for each test type. The tests were conducted on concretes allowed to cure for 7 days at prescribed room temperature for both unreinforced and reinforced PC and 28 days age for OCC.

**Figure 4.** Temperatures control through installation of thermocouples.

and reinforced PC and 28 days age for OCC. The setting and recorded temperatures were close and only a difference not exceed-The setting and recorded temperatures were close and only a difference not exceeding ±3 °C was reported (Figure 5). The setting and recorded temperatures were close and only a difference not exceeding ±3 ◦C was reported (Figure 5).

ing ±3 °C was reported (Figure 5).

**Figure 5.** Real time-Experimental temperature records vs. heating time. **Figure 5.** Real time-Experimental temperature records vs. heating time.

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

process and thereby in concrete bulk.

### *3.1. Fibers Adding Effects on the Density and Porosity Properties of PC*

*3.1. Fibers Adding Effects on the Density and Porosity Properties of PC* The effect of addition of short fibers on the density PC mixed is shown in Table 4. The results reveals that the inclusion of short fibers have a slight effect on the density of PC concrete. The density of the studied mixes ranged from 2.16 g/cm3 to 2.23 g/cm3 . In fact, a marginal decrease was observed for the entire mixes. The same trend can be observed for the carbon group of mixes as the density decreased with higher fiber dosages. The maximum difference in densities recorded over the mixes was 3.2%, which shows that incorporation of fibers with small quantities did not significantly affect the concrete density. In addition, it was noted that the density of polypropylene fibers is lower than that of polymer. This could be due to the deficiency of uniform distribution during the mixing The effect of addition of short fibers on the density PC mixed is shown in Table 4. The results reveals that the inclusion of short fibers have a slight effect on the density of PC concrete. The density of the studied mixes ranged from 2.16 g/cm3 to 2.23 g/cm3 . In fact, a marginal decrease was observed for the entire mixes. The same trend can be observed for the carbon group of mixes as the density decreased with higher fiber dosages. The maximum difference in densities recorded over the mixes was 3.2%, which shows that incorporation of fibers with small quantities did not significantly affect the concrete density. In addition, it was noted that the density of polypropylene fibers is lower than that of polymer. This could be due to the deficiency of uniform distribution during the mixing process and thereby in concrete bulk.

**Table 4.** Density and porous structure properties of studied concretes. **Table 4.** Density and porous structure properties of studied concretes.

The pore structures (i.e., mean pore size, pore size distribution, and various pores size proportions) were characterized by means of MPI and results are presented in Table 4. The introduction of fibers leads also to a decrease in the mean most distributed pore diameter dc. The increase in the porosity of PC reinforced with polypropylene fibers can be attributed to volume expansion of the fibers during mixing which conducts to an additional porosity. In general, the addition of short fibers to polymer concrete generates difficulties of the solid particles compaction causing an increase in the total porosity and a decrease in the most distributed pore diameter, dc as shown in Figure 6. Similar phenomena have been observed by Bentur et al. [41] and Barbuta et al. [42]. Although, the porosity of PC mixes increases with fiber content, but is still significantly lower than that of OCC samples. The pore structures (i.e., mean pore size, pore size distribution, and various pores size proportions) were characterized by means of MPI and results are presented in Table 4. The introduction of fibers leads also to a decrease in the mean most distributed pore diameter dc. The increase in the porosity of PC reinforced with polypropylene fibers can be attributed to volume expansion of the fibers during mixing which conducts to an additional porosity. In general, the addition of short fibers to polymer concrete generates difficulties of the solid particles compaction causing an increase in the total porosity and a decrease in the most distributed pore diameter, dc as shown in Figure 6. Similar phenomena have been observed by Bentur et al. [41] and Barbuta et al. [42]. Although, the porosity of PC mixes increases with fiber content, but is still significantly lower than that of OCC samples.

#### *3.2. Effects of Adding Fibers on the Modulus of Elasticity of PC Systems 3.2. Effects of Adding Fibers on the Modulus of Elasticity of PC Systems*

The dynamic Young's modulus Ed and the shear modulus are determined using UPV. The two dynamic modulus are given by the following Equations (5) and (6): The dynamic Young's modulus Ed and the shear modulus µd are determined using UPV. The two dynamic modulus are given by the following Equations (5) and (6):

$$\mathbf{V\_L} = \left(\frac{\mathbf{E(1-\vartheta)}}{\rho(1+\vartheta)(1-2\vartheta)}\right)^{\frac{1}{2}}\tag{5}$$

$$
\mu\_{\rm d} = \rho \left| \mathbf{V}\_{\rm T}^{2} \right. \tag{6}
$$

The value of Poisson's ratio is given by Equation (7): The value of Poisson's ratio is given by Equation (7):

$$\theta = \frac{1 - 2(\mathbf{V\_T/V\_L})^2}{2 - 2(\mathbf{V\_T/V\_L})^2} \tag{7}$$

In the Equation (5) to Equtaion (6), Ed represents the dynamic elasticity modulus (MPa), VL is the compressive P-wave velocity, *VT* is shear wave velocity (km/s), *ν* is Poisson's ratio, and is the density (kg/m3 ). In the Equation (5) to Equtaion (6), Ed represents the dynamic elasticity modulus (MPa), VL is the compressive P-wave velocity, VT is shear wave velocity (km/s), ν is Poisson's ratio, and *ρ* is the density (kg/m3 ).

The results of the experimental program revealed that the static elastic modulus (Es) decreased slightly as the fiber content increased (Figure 7a,b). This marginal variation is the result of two balanced phenomena: an increase due to the high elastic modulus of the carbon fibers and decrease due to the increase in porosity and the anisotropic feature. By increasing the fraction of fibers, the growth of porosity outweighs the contribution of fibers. The lowest rigidity is observed for PC-P1 because polypropylene fibers have a low elastic modulus similar to that of the epoxy binder. Moreover, PC-P1 has higher values of the total porosity compared to control sample. In the same context, Reis and Kumar[10,43] in their research paper confirmed that adding carbon and glass fibers do not improve the compressive elastic modulus of the composites but in opposite slight decrease was observed while carbon fiber reinforcements are incorporated. The results of the experimental program revealed that the static elastic modulus (Es) decreased slightly as the fiber content increased (Figure 7a,b). This marginal variation is the result of two balanced phenomena: an increase due to the high elastic modulus of the carbon fibers and decrease due to the increase in porosity and the anisotropic feature. By increasing the fraction of fibers, the growth of porosity outweighs the contribution of fibers. The lowest rigidity is observed for PC-P1 because polypropylene fibers have a low elastic modulus similar to that of the epoxy binder. Moreover, PC-P1 has higher values of the total porosity compared to control sample. In the same context, Reis and Kumar [10,43] in their research paper confirmed that adding carbon and glass fibers do not improve the compressive elastic modulus of the composites but in opposite slight decrease was observed while carbon fiber reinforcements are incorporated.

**Figure 7.** Elastic characteristics of polymer concretes: (**a**) Shear Modulus, (**b**) Young Modulus. **Figure 7.** Elastic characteristics of polymer concretes: (**a**) Shear Modulus, (**b**) Young Modulus.

By adding fibers to PC, all mixes showed indistinctly an increase in the modulus with sample age and being stabilized after a period of 5 to 7 days (Figure 8). This can be attributed to the progress of the curing reaction taken end at the 7 days as maximum leading to the formation of polymeric epoxy structures [16]. It should be mentioned that the modulus of elasticity records for unreinforced PC displayed are in accordance with the values described in literature for epoxy-based PC [11,44]. By adding fibers to PC, all mixes showed indistinctly an increase in the modulus with sample age and being stabilized after a period of 5 to 7 days (Figure 8). This can be attributed to the progress of the curing reaction taken end at the 7 days as maximum leading to the formation of polymeric epoxy structures [16]. It should be mentioned that the modulus of elasticity records for unreinforced PC displayed are in accordance with the values described in literature for epoxy-based PC [11,44]. *Polymers* **2023**, *15*, x 10 of 16

**Figure 8.** Elastic modulus vs. age for PC mixes.

[45].

pressive strength.

PC PC-C1 PC-C2 PC-P1

**PC Mixes**

**Bending strength (MPa)**

**Figure 8.** Elastic modulus vs. age for PC mixes.

*3.3. Fibers Adding Effects on the Mechanical Properties of PC*

25.9% for the same fiber content, comparing to PC-C1.

(**a**) (**b**)

0

20

40

**Compressive strength (MPa)**

60

80

100

reinforced with polypropylene and carbon fibers, regardless of the type of reinforcements. In fact, it was reported that polypropylene fibers as example show weak binder–fiber contacts and this decrease the performance of the final product, regardless of the binder type

However, the results show that flexural strength of PC made with carbon fibers increased with the increase in fibers content (Figure 9a). In average, flexural strengths increased 20% approximately for 2% fiber content, and compressive strength decreased

**Figure 9.** Bending and compressive resistances of different PC sets: (**a**) Flexural strength, (**b**) Com-

PC PC-C1 PC-C2 PC-P1 **PC Mixes**

12

17

22

**Young Modulus E (GPa).**

27

32

37

PC PC-C1 PC-C2 PC-P1

#### *3.3. Fibers Adding Effects on the Mechanical Properties of PC* tacts and this decrease the performance of the final product, regardless of the binder type [45].

*3.3. Fibers Adding Effects on the Mechanical Properties of PC*

2 4 6 9 12

**Age (days)**

**Figure 8.** Elastic modulus vs. age for PC mixes.

*Polymers* **2023**, *15*, x 10 of 16

The results of flexural and compressive strength of PC mixes are presented in Figure 9a,b. The composite systems exhibit lower flexural and compressive characteristics when reinforced with polypropylene and carbon fibers, regardless of the type of reinforcements. In fact, it was reported that polypropylene fibers as example show weak binder–fiber contacts and this decrease the performance of the final product, regardless of the binder type [45]. However, the results show that flexural strength of PC made with carbon fibers increased with the increase in fibers content (Figure 9a). In average, flexural strengths increased 20% approximately for 2% fiber content, and compressive strength decreased 25.9% for the same fiber content, comparing to PC-C1.

The results of flexural and compressive strength of PC mixes are presented in Figure 9a,b. The composite systems exhibit lower flexural and compressive characteristics when reinforced with polypropylene and carbon fibers, regardless of the type of reinforcements. In fact, it was reported that polypropylene fibers as example show weak binder–fiber con-

**Figure 9.** Bending and compressive resistances of different PC sets: (**a**) Flexural strength, (**b**) Compressive strength. **Figure 9.** Bending and compressive resistances of different PC sets: (**a**) Flexural strength, (**b**) Compressive strength.

However, the results show that flexural strength of PC made with carbon fibers increased with the increase in fibers content (Figure 9a). In average, flexural strengths increased 20% approximately for 2% fiber content, and compressive strength decreased 25.9% for the same fiber content, comparing to PC-C1.

The decrease in the mechanical resistances is probably due to poor adhesion between the fibers and the matrix and to the high values of porosity. This decrease is a sign of decrease in transverse bonds in polymer, which decrease the stiffness and increase ductility of polymer concrete [46]. In fact, adding 1% of carbon fibers does not affect the compressive strength and the elastic characteristics but decreases the flexure strength. In other words, fibers do not increase PC strength but their addition to the mixture diminished the signs of brittleness behavior of unreinforced polymer concrete [20], which is translated by a change in bending behavior from quasi-linear brittle to nonlinear ductile, a nonlinearity heightened by increasing the fiber content.

It can be concluded that the addition of fibers did not accomplish the expected reinforce or at least has the same strength characteristics as unreinforced PC [17]. Hence, only the toughness of polymer concretes reinforced with carbon fibers will be considered thereafter.

### *3.4. Fibers Adding Effects on Toughness at High Temperatures*

Figure 10a,b represent, respectively, the variation of stress intensity factor KIC and the fracture energy GF as a function of the exposure temperature for both the PC system and the OCC. It can be seen that PC is more resistant to crack propagation than OCC. By increasing the exposure temperature, the toughness is improved up to 225 ◦C for both concrete before declining slightly at higher temperatures (250 ◦C). Such a variation has been similarly observed for the flexural strengths mainly in the case of plain PC as a result of action of two coexistent and competitive phenomena: (i) post crosslinking initiated by heat and (ii) thermo-oxidative degradation of the epoxy polymer [47]. This variation is emphasized by a loss of bond between aggregates and the binder as demonstrated in a previous work [14].

The decrease in the mechanical resistances is probably due to poor adhesion between the fibers and the matrix and to the high values of porosity. This decrease is a sign of decrease in transverse bonds in polymer, which decrease the stiffness and increase ductility of polymer concrete [46]. In fact, adding 1% of carbon fibers does not affect the compressive strength and the elastic characteristics but decreases the flexure strength. In other words, fibers do not increase PC strength but their addition to the mixture diminished the signs of brittleness behavior of unreinforced polymer concrete [20], which is translated by a change in bending behavior from quasi-linear brittle to nonlinear ductile, a nonlinearity

It can be concluded that the addition of fibers did not accomplish the expected reinforce or at least has the same strength characteristics as unreinforced PC [17]. Hence, only the toughness of polymer concretes reinforced with carbon fibers will be considered there-

Figure 10a,b represent, respectively, the variation of stress intensity factor KIC and the fracture energy GF as a function of the exposure temperature for both the PC system and the OCC. It can be seen that PC is more resistant to crack propagation than OCC. By increasing the exposure temperature, the toughness is improved up to 225 °C for both concrete before declining slightly at higher temperatures (250 °C). Such a variation has been similarly observed for the flexural strengths mainly in the case of plain PC as a result of action of two coexistent and competitive phenomena: (i) post crosslinking initiated by heat and (ii) thermo-oxidative degradation of the epoxy polymer [47]. This variation is emphasized by a loss of bond between aggregates and the binder as demonstrated in a

heightened by increasing the fiber content.

*3.4. Fibers Adding Effects on Toughness at High Temperatures*

after.

previous work [14].

The effect of carbon fiber introduction on the fracture properties of PC is studied and the results are depicted on Figure 10a,b. It can be seen that adding 1% of carbon fibers to polymer concretes results in an enhancement on its fracture properties at room temperature, but this effect is vanished by increasing the exposure temperature as a result of the decrease in bond strength in the fiber–matrix interface while rising temperature [14] added to the other types of degradation mentioned earlier. The effect of carbon fiber introduction on the fracture properties of PC is studied and the results are depicted on Figure 10a,b. It can be seen that adding 1% of carbon fibers to polymer concretes results in an enhancement on its fracture properties at room temperature, but this effect is vanished by increasing the exposure temperature as a result of the decrease in bond strength in the fiber–matrix interface while rising temperature [14] added to the other types of degradation mentioned earlier.

Figure 11 shows the load–CMOD graph for the tested notched beams as a function of heating temperature. The overall load–CMOD curves have presented a smooth softening curves after the peak loads which indicates that tests were conducted under a stable test regime. All curves observed show a similar general trend for all specimens. Figure 11 shows the load–CMOD graph for the tested notched beams as a function of heating temperature. The overall load–CMOD curves have presented a smooth softening curves after the peak loads which indicates that tests were conducted under a stable test regime. All curves observed show a similar general trend for all specimens.