Composites Part B 247 (2022) 110342
Available online 7 October 2022
1359-8368/© 2022 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by- nc-nd/4.0/).
Ultrahigh strength and modulus of polyimide-carbon nanotube based carbon and graphitic fibers with superior electrical and thermal conductivities for advanced composite applications
Seo Gyun Kim
a,1, So Jeong Heo
a,b,1, Sungyong Kim
a,c, Junghwan Kim
a, Sang One Kim
a,c, Dongju Lee
a,d, Suhun Lee
a, Jungwon Kim
a, Nam-Ho You
a, Minkook Kim
a, Hwan Chul Kim
c, Han Gi Chae
b,**, Bon-Cheol Ku
a,e,*aInstitute of Advanced Composite Materials, Korea Institute of Science and Technology (KIST), Wanju, 55324, Republic of Korea
bDepartment of Materials Science and Engineering, Ulsan National Institute of Science and Technology (UNIST), Ulsan, 44919, Republic of Korea
cDepartment of Carbon Materials and Fiber Engineering, Jeonbuk National University, Jeonju, 54896, Republic of Korea
dGraduate School of Convergence Science and Technology, Seoul National University, Gyeonggi-do, 16229, Republic of Korea
eDepartment of Nano Convergence, Jeonbuk National University, Jeonju, 54896, Republic of Korea
A R T I C L E I N F O Keywords:
Carbon nanotubes Polyimide Carbon fibers
Carbon-fiber-reinforced polymer composites Mechanical properties
A B S T R A C T
Development of carbon fibers (CFs) with high strength and high modulus for structural applications in CF- reinforced polymer (CFRP) industry has been a challenge. Herein, we propose a method for manufacturing highly oriented polymer–carbon nanotube (CNT) composite fibers having high strength (4.8 ± 0.2 GPa), modulus (390 ±48 GPa), and electrical conductivity (5.75 ±0.84 MS m-1) by a liquid crystalline wet-spinning process. The use of chlorosulfonic acid (CSA) as a solvent for CNTs and polyimide (PI) promotes dispersion and enables the production of high-performance composite fibers. In addition, the functional groups of PI in com- posite fibers improve the interfacial shear strength with epoxy resin without sizing additives by 72% compared to that of CNT fibers. Carbonization and graphitization of the composite fibers with an optimal ratio of PI (30%) and CNT cause significant improvement in their mechanical (tensile strength; 6.21 ±0.3 GPa and modulus; 701 ± 47 GPa) and thermal properties (496 ±38 W m−1 K−1) by reducing voids and improving orientation. We believe that the polymer–CNT composites and their CFs with high strength and high modulus would be the next- generation CFs for aerospace and defense industry.
1. Introduction
Carbon fibers (CFs) have been predominantly used as reinforcing materials of composites. Among various applications, CFs have attracted significant attention in aerospace and defense industry because of their exceptional mechanical properties [1–4]. In regard to their tensile properties, polyacrylonitrile (PAN)-based CFs with high strength and pitch-based CFs with high modulus are well known in CF industry. For example, T1100G (PAN-based CF, Toray) with a tensile strength of 7.0 GPa has relatively low tensile modulus (324 GPa), and K13D2U (pitch-based CF, Mitsubishi) with a tensile modulus of 935 GPa has
relatively low tensile strength (3.7 GPa) [5,6]. Thus, developing CFs exhibiting both high strength and high modulus has been a challenge [7]. Even though recent studies have demonstrated that incorporation of CNTs improve the mechanical properties of PAN-based CFs, issues such as poor dispersion and limited amount of CNTs in polymer matrix still ought to be addressed [7–12].
More recent approaches utilizing pure CNT fibers by wet-spinning exhibited concurrent improvement in tensile strength and tensile modulus [6,13–18]. Recently, we successfully processed CNT-based graphitic fibers with strength and modulus of 6.5 GPa and 600 GPa, respectively [16]. These fibers were manufactured by wet-spinning an
* Corresponding author. Institute of Advanced Composite Materials, Korea Institute of Science and Technology (KIST), Wanju, 55324, Republic of Korea.
** Corresponding author. Department of Materials Science and Engineering, Ulsan National Institute of Science and Technology (UNIST), Ulsan, 44919, Republic of Korea.
E-mail addresses: [email protected] (H.G. Chae), [email protected] (B.-C. Ku).
1 S.G. Kim and S.J. Heo contributed equally to this work.
Contents lists available at ScienceDirect
Composites Part B
journal homepage: www.elsevier.com/locate/compositesb
https://doi.org/10.1016/j.compositesb.2022.110342
Received 29 June 2022; Received in revised form 16 September 2022; Accepted 4 October 2022
anisotropic liquid-crystalline (LC) phase of CNT solution using chlor- osulfonic acid (CSA). Herein, CSA-based LC solutions of CNTs enable extremely high orientation and packing density of CNT fibers [14–17].
However, some studies are being conducted to manufacture CNT fibers with mild acids [19] because CSA is a strong acid and challenging to use.
In addition, the cost of CNTs with high purity and crystallinity is high ($2,000 to $100,000/kg), which further makes commercialization of the process challenging [6]. CSA, which is a solvent for CNTs, can be also considered as a solvent for various polymers because of its extremely high acidity (Hammett acidity function, H0 =13.8) [20,21]. It readily protonates aromatic hydrocarbons, thereby promoting the dissolution of various polymers containing aromatic groups. In addition, CSA acts as a polar solvent and dissolves the polymers that contain polar functional groups, such as alcohols, carboxylic acids, and amines, in their proton- ated form. Therefore, CSA can be an effective solvent in manufacturing polymer–CNT composite fibers with high CNT content. CNTs and the polymer can be simply and effectively dispersed and hybridized without physical or chemical treatment, and high-performance fibers can be obtained only by wet-spinning without high-temperature heat treat- ment. Moreover, unlike CF having a low surface energy, a polymer composite containing various functional groups can increase the inter- action with a resin such as epoxy by imparting functionality to the fiber and can expect the sizing effect of fibers for carbon fiber reinforced polymers (CFRPs).
In this study, we propose a novel process to fabricate wet-spun pol- yimide (PI)–CNT composite fibers and their CFs with high strength and high modulus. The PI–CNT fibers by LC spinning have high mechanical and electrical conductivity. The composite fibers showed improved interfacial shear strength for the epoxy. The proposed manufacturing process is expected to result in the manufacture of composite fibers and CFs with high strength and modulus.
2. Experiments 2.1. Materials
Single-walled CNTs were procured from Meijo Nano Carbon (DX-2).
CSA (Sigma-Aldrich) was used as the solvent for the CNTs and polyamic acid (PAA, SKC Kolon PI). The molecular structure of PAA was Poly (pyromellitic dianhydride-co-4,4′-oxydianiline) (PMDA-ODA) as shown in Fig. S1. Acetone (Daejung, 99.5%) was used as the coagulant during wet-spinning.
2.2. Fabrication of PI–CNT composite fibers
The CNT/PAA/CSA solution was spun through a single needle or multinozzle with an inner diameter of 0.11–0.26 mm. The flow rate was controlled in the range of 0.05–0.1 mL min−1 using a syringe pump (Fusion 710, Chemyx). The draw ratios above 2.0 were controlled by adjusting the winding rate at a fixed flow rate (0.1 mL min−1). The bath length for coagulation was 40 cm. The CNT/PAA fibers coagulated by acetone were washed with distilled water for 2 h and dried at 70 ◦C for 24 h in a vacuum oven. The composite fibers were imidized by contin- uously heat treating it for 0.5 h at 80 ◦C, 1 h at 175 ◦C, 1 h at 235 ◦C and 1 h at 350 ◦C under nitrogen atmosphere; the temperature at each stage was increased at a rate of 10 ◦C min−1.
2.3. Interfacial shear strength
The IFSS of the PI–CNT fibers in the epoxy resin was measured using a microdroplet test (Fig. 5a). Epoxy resin (YD-128, Kukdo) and a curing agent (MDA-60, Kukdo) were mixed in a weight ratio of 100:30, and the embedded length of microdroplets was 200–250 μm. The PI–CNT fibers were pulled out at a speed of 0.1 mm min−1 using a 2.5 N load cell and a universal testing machine (Instron 4464, Instron). More than 25 indi- vidual tests were conducted for each IFSS measurement.
2.4. Heat treatment of composite fibers
The composite fibers were fixed on a graphite sheet under constant weight. For this, a constant load (0.031 N tex−1) was applied to each precursor fiber by using clips weighing 0.754 g. The samples were heat treated at 1,400 ◦C and 2,700 ◦C using a tube-type carbonization furnace and graphitization furnace, respectively. The temperature was increased at 10 ◦C min−1 up to 1,400 ◦C and 2 ◦C min−1 up to 2,700 ◦C under an Ar atmosphere, and each temperature was maintained for 30 min.
2.5. Measurements of mechanical properties
The tensile strength and tensile modulus were measured using FAVIMAT+(Textechno), and the fibers with a length of 25 mm were used for the tensile test. The elongation rate was 2 mm min−1, and the pre-tension force was 1–2 cN. The specific strengths were measured at least 30 times.
2.6. Linear density and specific density of fibers
The linear density was measured by calculating the resonance fre- quency obtained using a vibroscope of FAVIMAT+(Textechno) [22].
The resonance frequency of the fibers was measured while increasing the pre-tension from 1 cN to 2 cN at a rate of 2 mm min−1. The measured linear density was also reconfirmed by the weight and length method as shown in Fig. S2. The weight per unit length of fiber having 0.17 tex (PI 30% composite fiber) was measured, and the average linear density (tex) was in good agreement with that measured from FAVIMAT+. The spe- cific density of the fibers was measured by a density gradient column (POLYTEST, Ray-Ran, UK). Two liquids (benzene and 1,1,2,2,-tetrabro- moethane) were used to build the density gradient, and glass beads with accurately known density floated in the column [23]. The fibers were inserted and left for 24 h in the column. The density of the fibers measured by the density gradient column method was reconfirmed through the linear density (tex) obtained from FAVIMAT+ and the cross-sectional area obtained from SEM.
2.7. Measurements of electrical and thermal properties
The electrical conductivity of the CNT fibers was measured using the four-point probe method with a probe station (MST-4000A, MS Tech).
The thermal conductivity of the PI–CNT fibers was measured using the self-heating method [16,17,24]. For this, the heat generated due to Joule heating by passing direct current through the fiber and the resulting temperature rise of the fiber were measured by the thermos-resistance property of each fiber with a four-point measure- ment. Heat transfer through convection and radiation were negligible due to high vacuum (~10−6 Torr) and the temperature rise was small (below 10 K). The heat generated by Joule heating was measured with a source meter (Keithley 6221) and a nanovoltmeter (Keithley 2182A).
The length and cross-sectional area of the fiber were measured by optical microscopy and SEM.
2.8. Characterization
Information about internal voids and crystal structure was obtained from 2D small-angle X-ray scattering (SAXS) and wide-angle X-ray scattering (WAXS) patterns using Rigaku Micromax-007 (operated at 45 kV, 66 mA, λ =0.154 nm) equipped with a Rigaku R-axis IV++detec- tion system at different sample-to-detector distances (SDDs). The SDDs for SAXS and WAXS were 3.06 m and 0.25 m, respectively. Information about the (002) crystal plane was obtained from the 2D WAXS pattern using IGOR Pro. The orientation factor of the (002) plane was obtained using an azimuthal scan at 2θ ~ 26◦using Hermans’ orientation factor and Wilchinsky’s equation. The in-plane and out-of-plane crystallite sizes were determined using the Scherrer equation with K =0.9. The
cross sections of the fibers were prepared using the focused ion beam of an FEI-Helios scanning electron microscope (SEM) and observed using high-resolution transmission electron microscopy (HR-TEM, FEI-Titan Cubed 60–300). X-ray photoelectron spectroscopy (XPS) spectra were analyzed by Thermo Scientific K-alpha (Thermo VG, USA) with mono- chromatic Al Kα (1486.6 eV). The functional groups of PAA and PI were characterized by Fourier transform infrared (FT-IR) spectroscopy (Nicolet IS10, USA) with 256 scans (4 cm−1 resolution) in the ATR mode. The orientation factor of pristine CNT fiber was confirmed by polarized Raman spectroscopy (InVia Reflex, Renishaw) with 514 nm excitation wavelengths [14,15]. The liquid crystal (LC) phase of the CNT solutions was obtained using an optical microscope with a cross-polarizing filter (L150, Nikon).
3. Results and discussion 3.1. Fabrication of composite fibers
CSA protonates CNTs so that they exist individually and spontane- ously form LCs [17–20]. The excellent dispersibility of CNTs in CSA allows for the fabrication of high-performance CNT fibers [6,13–20]. PI has excellent mechanical properties and thermal stability and, hence, is widely used for manufacturing composites [25,26]. However, PI has poor solubility in most solvents including CSA. Thus, it is difficult to obtain composite fibers with PI through wet-spinning. The solubility of PAA, the precursor of PI is better than that of PI because PAA has (i) a smaller rigid and flat aromatic surface and (ii) more polar functional
groups that promote dissolution in polar solvents.
In this study, PI–CNT composite fibers were fabricated by dispersing PAA in a CNT/CSA solution beforehand and imidizing it to PI. CSA was used as a solvent to facilitate the dispersion of CNTs and manufacture PI–CNT composite fibers with a high CNT content. However, as the amide linkages in the polymer backbone are vulnerable to hydrolysis under highly acidic conditions, the molecular weight of PI can decrease.
Therefore, after dissolving PAA in CSA and stirring the solution for 12–96 h (Fig. 1a), the precipitated PAA powder was dissolved again in N-methyl-2-pyrrolidone (NMP) for viscosity and molecular weight measurements (Fig. 1b). The specific viscosity (ηsp) and intrinsic vis- cosity ([η]) of the PAA/NMP solution are shown in Fig. 1a and b, respectively. ηsp decreased until 48 h and remained almost constant thereafter. [η]can be obtained from concentration (c)-dependent ηsp, as follows [27]:
[η] =lim
c→0
(ηsp c
) (1)
The molecular weight of PAA can be estimated using the Mark–Hauwink equation ([η] =KMaw, K=2.3×10−4, a =0.78, and Mw
is the molecular weight) [28]. The molecular weight of PAA before dissolution in CSA was 98,120 g mol−1 ([η] = 1.75dL/g), which decreased by 47%–52,190 g mol−1 ([η] =1.10dL/g)after 48 h. In this study, a CNT/PAA solution immersed in CSA for 72 h was used to eliminate the effect of the molecular weight reduction (Fig. 1a).
The IR spectra confirmed that intrinsic functional groups of PAA and PI were present (Fig. S3). In the IR spectrum, the structural differences
Fig. 1. Fabrication of PI–CNT composite fibers. (a) Specific viscosity of PAA/NMP solution as a function of duration of immersion of PAA in CSA. (b) Estimation of intrinsic viscosity of PAA/NMP solution before and after PAA is immersed in CSA. (c) Polarized optical images of CNT/CSA and CNT/PAA/CSA solutions with a concentration of 10 mg mL−1 for observing LC phase. Scale bar is 200 μm. (d) Wet-spinning process for fabricating composite fibers. (e) SEM images of twisted and knotted composite fibers with PI 10%. (f) SEM images of cross-sections of CNT and composite fibers.
between PAA and PI were indicated through the bands of different functional groups. Because PAA had amide groups in its structure, C––O stretch bands (1,721 cm−1, 1,650 cm−1) and C–N stretch bands (1,542 cm−1) were observed in PAA spectra, while the bands about amide group were not in PI spectra. Therefore, C––O stretching was observed at 1,778 cm−1 and 1,716 cm−1 which were different from C––O stretching of PAA. Additionally, stretching of C–N–C (1,375 cm−1) and bending of C––O (725 cm−1) were observed in the IR spectrum of PI (Fig. S3). Due to CSA as role of sulfonating agent for polymer, the functional groups of PAA and PI were also measured through XPS. The XPS and FT-IR results confirmed that PAA and PI were partially sulfonated and the sulfonated group was removed after carbonization process (Figs. S3 and S4).
CNT/CSA solutions are well known as lyotropic LCs [17,20]. The LC behavior of CNT solutions can be quite useful for determining the orientation and packing of CNTs during fiber spinning [17]. In addition
to the CNT/CSA solution, the LC behavior of the solutions containing PAA were studied through polarized optical imaging (Fig. 1c). The LC state was observed even at a PAA content of 60%. Long CNTs (~10 μ m) with a high aspect ratio [17], which were relatively more macromo- lecular than PAA, had a significant effect on LC formation in CSA. The spinning of LC solutions in the nematic phase is a prerequisite for improving the alignment and packing density of fibers [13–17]. In this study, composite fibers were prepared through wet-spinning using 10 mg mL−1 of the LC solution (Fig. 1d). Spinning consists of a drawing process, which applies a stretching force. This is the most important step for improving the molecular orientation [17]. In addition, it consists of a coagulation process, which is one of the unit operations used to control macrovoids inside the fibers [17]. A draw ratio of 2.0 or higher was optimized in a stable operation window based on the theoretical HSL model of wet-spinning [15]. Additionally, acetone was used as a
Fig. 2. Mechanical, electrical, and thermal properties of the fibers. Stress–strain curves of (a) PI–CNT composite fibers, (b) carbon fibers heat-treated at 1,400 ◦C, and (c) graphitic fibers heat-treated at 2,700 ◦C. (d) Tensile strength, (e) tensile modulus, (f) toughness, and (g) electrical conductivity of the fibers depending on the PI content. (h) Comparison of PI–CNT fibers with other composite fibers and CFs. (i) Radar charts of composite, carbon, and graphitic fibers of PI–CNT and commercial carbon fibers. Each of the properties were normalized to the highest value of the fibers.
coagulant to determine the optimal macroscopic properties [29]. A large-scale continuous spinning process was performed, and the diam- eter of the fibers was controlled using the nozzle diameter or number of holes in the nozzle. The manufactured PI–CNT composite fibers through imidization exhibited excellent flexibility for twisting and knotting (Fig. 1e).
3.2. Macroscopic properties
The PI–CNT composite fibers had a high tensile strength and tensile modulus, which were comparable to those of conventional PAN-based or pitch-based CFs (Fig. 2). The tensile strength of the PI–CNT composite fibers (3.4–4.8 GPa) without high-temperature treatment was compa- rable to that of CF (T700G, Toray), and its tensile modulus was 40%– 63% higher than that of T700G. The PI–CNT composite fibers were carbonized and graphitized by heat treatment at 1,400 ◦C and 2,700 ◦C, respectively, and SEM images of cross-sections are shown in Fig. 1f. The heat treatment of the composite fibers caused significant changes in the mechanical and electrical properties of the fibers (Fig. 2a–h). The tensile strength was observed to be maximum at 30% PI regardless of the heat- treatment temperature, and the fiber with 30% PI heat-treated at 1,400 ◦C also had the highest strength (6.21 ±0.3 GPa), as shown in Fig. 2d. The tensile strengths of the PI–CNT-based CFs heat treated at 1,400 ◦C (3.79–6.2 GPa) were similar to that of CF (IM9, Hexcel), and their modulus (357–528 GPa) was 32%–68% higher than that of PAN- based CFs. The tensile modulus was observed to be the highest for the graphitized fibers treated at 2,700 ◦C, which decreased with increasing PI content (Fig. 2e). The fibers that were heat treated at 2,700 ◦C were comparable to pitch-based CFs. They showed an excellent tensile modulus of 406–841 GPa, and their tensile strength was 13%–30%
higher than that of pitch-based CFs. The mechanical properties of the PI–CNT composite, carbonized, and graphitized fibers were consider- ably superior to those of previously reported polymer–CNT-based fibers (Fig. 2h) [30–39]. Here, it is worth noting that the PI–CNT composite fibers exhibited both high tensile strength and tensile modulus compared to conventional CFs. In addition, by changing the heat-treatment conditions, fibers with a higher tensile strength or tensile modulus can be manufactured from the same precursor.
Compared to those of the heat-treated fibers, the electrical and thermal conductivities of the PI–CNT composite fibers were superior (Fig. 2g and S5). The electrical conductivity was observed to be 8–80 times higher than T1100G (0.071 MS m−1) and K13D2U (0.67 MS m−1).
Apart from that, the PI–CNT based composite, carbon, and graphitic fibers have excellent thermal conductivity of 335–496 W m−1 K−1, which is comparable to that of pitch-based fibers (Fig. S5). Such excel- lent properties of PI–CNT-based fibers make them superior to conven- tional CFs (Fig. 2i).
3.3. Microstructure
3.3.1. Microstructures of composite fibers
The excellent tensile properties of the fabricated PI–CNT fibers can be inferred from their microstructure (Fig. 3). The SAXS 2D patterns (Fig. 3a) show that the microstructure changed with the PI content. The tensile properties were strongly affected by the void structure and void orientation, as can be observed from the 2D SAXS patterns. In the 2D SAXS patterns (Fig. 3a), sharp edges became round with an increase in the PI content, indicating an increase in void misalignment along the fiber axis (Fig. S6) with an increase in PI content and its bulk phase. The 2D SAXS pattern was also used to estimate the orientation of microvoids
Fig. 3. Microstructures of PI–CNT composite fibers. (a) 2D SAXS patterns of PI–CNT fibers with 0%, 30%, and 60% PI. (b) Volume and misalignment of voids vs. PI content. (c) Relationship between void volume and tensile strength. (d) Orientation factor of (002) plane vs. PI content. (e) Schematic illustration of CNT orientation during wet-spinning process.
along the fiber axis in azimuthal scans of the fiber’s equator. Next, using Ruland’s method and classical Porod’s law applying fractal theory, the respective void size and misalignment were calculated [40,41]. The void volume was calculated by assuming cylindrical voids. The PI–CNT fibers with 30% PI had a relatively small void volume (0.11 nm3) and a low degree of void misalignment along the fiber axis (0.35 radians) (Fig. 3b and c).
Hermans’ orientation factor of the (002) carbon plane was calculated using Wilchinsky’s equation [42]. The orientation factor of pristine CNT fiber measured by polarized Raman spectroscopy was 0.89–0.91, which was similar to that obtained from WAXS (Fig. S7) [15]. The (002) plane of the PI–CNT fibers exhibited better alignment than that of the pristine CNT fibers (Fig. 3d). In general, viscoelastic fluid under high shear generates normal stress, which induces die swell at the exit of spinning nozzle [43]. As a result, poor molecular orientation may be obtained.
However, proper PI content in the solution may prevent disorder of CNTs because they can act as a binder (Fig. 3e). In addition, the improvement in orientation leads to better molecular packing and to a reduction in the void volume. Eventually, these lead to improved tensile
properties. The electrical conductivity of the PI–CNT composite fiber increased from 5.18 MS m−1 to 5.75 MS m−1 up to 30% PI content (Fig. 2g). It can be observed that the orientation of the (002) plane and the void volume also affect the electrical conductivity. However, above 30% PI, the orientation factor of the (002) plane gradually decreased, the void volume increased, and misalignment of voids occurred. This suggests that PI molecules formed a bulk phase and were poorly aligned along the fiber axis. The highest tensile strength (4.8 ±0.2 GPa) and modulus (390 ±48 GPa) were obtained at 30% PI, which implied that the tensile strength, modulus, and electrical conductivity were functions of the void volume and alignment of the (002) plane (Fig. 3b–d). This suggests the existence of an optimal structural balance between PI and the CNTs in the composite fibers.
3.3.2. Microstructures of carbonized and graphitized fibers
From the XPS analysis, the carbon contents of the fibers heat-treated at 1,400 ◦C and 2,700 ◦C were 92 ± 1% and 97 ±1%, respectively (Fig. S8). The tensile properties of the precursor are crucial for deter- mining the tensile strength and modulus after carbonization and
Fig. 4. Correlation between tensile properties and microstructure of carbon and graphitic fibers of PI–CNT. (a) 2D SAXS patterns of PI–CNT fiber with 30% PI heat treated at 1,400 ◦C and 2,700 ◦C. (b) 2D WAXS patterns of PI–CNT fiber with 30% PI heat treated at 1,400 ◦C and 2,700 ◦C. (c) Relationship between tensile strength and void thickness. (d) Relationship between tensile modulus and orientation factor of (002) plane. (e) HR-TEM images of cross-sections of fibers heat-treated at 1,400 ◦C and 2,700 ◦C. (f) Schematic illustration of the cross-section of PI–CNT fibers as a function of the heat-treatment temperature.
graphitization [44]. The well-aligned structure of PI can be converted to CFs without conventional stabilization process for PAN precursor because of the aromatic ring on the backbone [45]. Thus, the heat-treatment process is simpler than that for PAN-based CFs. Fig. 4 shows the correlation between the microstructure and tensile properties of the heat-treated PI–CNT fibers. The tensile strength of the heat-treated PI–CNT fibers with 30% PI increased to 6.21 ±0.3 at 1, 400 ◦C and 4.54 ±0.4 GPa at 2,700 ◦C, and it decreased for PI content beyond 30% (Fig. 2d). The highest strength was derived from strong precursor fibers, and the variation in strength with the PI content was similar to that of the precursor. However, the variation of the modulus with the annealing temperature was different from that of strength, particularly for fibers annealed at 2,700 ◦C (Fig. 4c and d). This indicates that different aspects of the microstructure affected the strength and modulus of the fibers. SAXS and WAXS were used to investigate the effects of structural evolution on the fiber properties according to the heat-treatment temperature, a schematic illustration of which is pre- sented in Fig. 4f. As shown in Fig. 4a and b, the microstructures and nanostructures of the fibers varied with the annealing temperature. The equatorial streak of the fibers that were heat treated at 2,700 ◦C was longer than that of the fibers heat treated at 1,400 ◦C, indicating an improvement in the tensile modulus of the fibers due to the increased void orientation. In addition, the 2D SAXS pattern showed an ellipsoidal shape of the fibers that were heat treated at 2,700 ◦C (Fig. 4a and b), indicating a decrease in the tensile strength resulting from an increase in the void thickness. Changes in void was derived from variation of the nanostructure as shown in Fig. 4c [46]. At 1,400 ◦C, PI molecules were converted to a turbostratic structure, and the nongraphitized carbon between the turbostratic structures was additionally pyrolyzed at 2, 700 ◦C, resulting in an increase in the in-plane and out-of-plane sizes of the crystals, as shown in the TEM images of the carbonized and graph- itized fiber (Fig. 4e and S9). The resulting graphite-like layers led to an increase in the void thickness owing to folding, splitting, or tilting [47].
This deteriorated the tensile strength of the fibers, as reported in a previous study [48,49]. It is confirmed that the tensile strengths of the fibers heat-treated at 2,700 ◦C are lower than those of the fibers heat-treated at 1,400 ◦C at the same PI content (Fig. 4c).
2D WAXS patterns showed that the orientation factor of the (002) plane increased with the heat-treatment temperature (Fig. 4b). The orientation factor of fibers with 30% PI was the highest, which showed the highest tensile modulus among the fibers heat treated at 1,400 ◦C.
Similarly, the pristine CNT fiber, which had the highest orientation factor at 2,700 ◦C, also showed the highest tensile modulus. However, it decreased as the PI content increased at 2,700 ◦C (Fig. 4d). This indi- cated that graphite-like layers were stacked more irregularly as the PI content increased when less-ordered PI molecules were graphitized.
Unlike the tensile strength determined by voids, the tensile modulus of the PI–CNT fibers increased with the annealing temperature because of the well-aligned (002) plane and the increase in the crystal size (La, and Lc) along the in-plane and out-of-plane directions (Fig. 4d and S9). The PI–CNT fibers heat-treated at 2,700 ◦C exhibited the highest thermal
conductivity (496 ±38 W m−1 K−1), and this was attributed to a large longitudinal crystal size and a corresponding reduction in phonon scattering (Figs. S5 and S9) [16,50]. It has been reported that the development of La revealed a proportional relationship with thermal conductivity [16]. The electrical conductivity decreased with increasing annealing temperature (Fig. 2g). This is mainly due to the dedoping effect of CSA at 1,400 ◦C and fusing of unzipped carbon sheet at 2,700 ◦C [16]. The fibers annealed at 1,400 ◦C also exhibited lower electrical conductivity than the pristine fibers despite carbonization of the PI. The differences in the electrical conductivities of the CNT fiber and com- posite fibers with 30% PI before and after heat-treatment at 1,400 ◦C suggests that the doping effect is relatively large.
3.4. Interfacial shear strength
The IFSS between the fibers and matrix is one of the main factors that determines the mechanical properties of fiber-reinforced polymer (FRP) composites. The IFSS of the PI–CNT fibers with 30% PI was 16.2 MPa, which was 72% higher than that of the pristine CNT fibers (6.8 MPa).
However, the IFSS decreased to 11.9 MPa and 11.0 MPa at PI contents of 50% and 60%, respectively (Fig. 5c). The shear failure in the micro- droplet debonding test did not occur at the interface between the epoxy resin and CNT fibers, but the outermost CNT layer of the PI–CNT fibers peeled off. The exfoliated CNTs were bonded to the fractured surface of the droplets after the test (Fig. 5b). In other words, the IFSS of the PI–CNT fibers was determined not only by the interfacial bonding strength between the CNT fibers and epoxy resin but also the shear bonding strength between the CNTs. As a result, the behavior of the IFSS of the PI–CNT fibers was similar to that of their toughness and tensile strength. When the PI content was above 50%, the void fraction of the PI–CNT fibers increased significantly and the shear bond strength be- tween CNTs decreased. As the interfacial bonding force between the PI–CNT fibers and epoxy resin is stronger than the shear strength of the PI–CNT fibers, a complete reinforcing effect can be expected when epoxy-based FRP composites are constructed using PI–CNT fibers. The IFSS of the PI–CNT fibers heat treated at 1,400 ◦C decreased from 10.1 MPa, 16.2 MPa, and 11.9 MPa to 7.7 MPa, 9.0 MPa, and 8.9 MPa for PI contents of 10%, 30%, and 50%, respectively. The PI in composite fibers that have not been heat treated has more functional groups such as N or O, which can make the fiber less hydrophobic and have a higher surface energy than carbonized fibers (Figs. S3, S4, and S8) [51]. This, therefore, causes the interfacial strength between the fibers and resin to be stronger.
4. Conclusions
PI–CNT composite fibers with high strength, modulus, and conduc- tivities were successfully fabricated using CSA as a solvent to disperse CNT and polymer. Even without carbonization at high temperatures, strong and conductive composite fibers with tensile strength of 4.8 ± 0.2 GPa, tensile modulus of 390 ±48 GPa, and electrical conductivity of
Fig. 5.Interfacial shear strength (IFSS) of the PI–CNT fibers. (a) Embedded length of the microdroplet. (b) Droplet failure after the test. (c) IFSS of the CNT, PI–CNT fibers, and fibers heat treated at 1,400 ◦C.
5.75 ±0.84 MS m−1 were produced. The IFSS of the PI–CNT composite fibers increased by 72% owing to the sizing effect of the functional groups of PI. Additionally, the tensile properties of the fibers were significantly improved after heat treatment at the carbonization tem- perature. The carbonized fibers with 30% PI content showed excellent mechanical properties, with a tensile strength of 6.21 ± 0.3 GPa at 1,400 ◦C and a tensile modulus of 701 ±47 GPa at 2,700 ◦C. The thermal conductivity of the graphitized fiber at 2,700 ◦C exhibited excellent results of up to 496 ±38 W m−1 K−1. The optimal PI content improved the orientation of the fibers and reduced the void thickness, resulting in successful fabrication of high-performance composite fibers with a synergistic effect between polymer and CNT. The manufacturing of high-performance composites with a high CNT content is therefore expected to enable the fabrication of customized composite fibers with the structures and properties required for CFRPs by utilizing various polymers.
CRediT authorship contribution statement
Seo Gyun Kim: Investigation, Conceptualization, Methodology, Writing – original draft. So Jeong Heo: Investigation, Conceptualiza- tion, Methodology, Writing – original draft. Sungyong Kim: Formal analysis, Validation. Junghwan Kim: Formal analysis, Validation. Sang One Kim: Formal analysis, Validation. Dongju Lee: Formal analysis, Validation. Suhun Lee: Formal analysis, Validation. Jungwon Kim:
Formal analysis, Validation. Nam-Ho You: Formal analysis, Validation.
Minkook Kim: Formal analysis, Investigation, Validation. Hwan Chul Kim: Formal analysis, Validation. Han Gi Chae: Investigation, Conceptualization, Writing – review & editing, Supervision. Bon-Cheol Ku: Investigation, Conceptualization, Supervision, Writing – review &
editing, Funding acquisition.
Declaration of competing interest
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
Data availability
Data will be made available on request.
Acknowledgments
This work was supported by grants from the Open Research Program (2E31902) and K-Lab of the Korea Institute of Science and Technology (KIST) and the grant (20011577) from Ministry of Trade, Industry and Energy (MOTIE) in Korea. This work was also supported by the National Research Foundation of Korea (2021R1A2C200440411).
Appendix A. Supplementary data
Supplementary data to this article can be found online at https://doi.
org/10.1016/j.compositesb.2022.110342.
References
[1] Zheng H, Zhang W, Li B, Zhu J, Wang C, Song G, et al. Recent advances of interphases in carbon fiber-reinforced polymer composites: a review. Compos B Eng 2022;233:109639.
[2] Peijs T, Kirschbaum R, Lemstra PJ. Chapter 5: a critical review of carbon fiber and related products from an industrial perspective. Adv Ind Eng Polym Res 2022;5:
90–106.
[3] Cheon J, Kim M. Impact resistance and interlaminar shear strength enhancement of carbon fiber reinforced thermoplastic composites by introducing MWCNT- anchored carbon fiber. Compos B Eng 2021;217:108872.
[4] Sayam A, Rahman ANMM, Rahman MdS, Smriti SA, Ahmed F, Rabbi MdF, et al.
A review on carbon fiber-reinforced hierarchical composites: mechanical
performance, manufacturing process, structural applications and allied challenges.
Carbon Lett 2022. https://doi.org/10.1007/s42823-022-00358-2.
[5] Chae HG, Newcomb BA, Gulgunje PV, Liu Y, Gupta KK, Kamath MG, et al. High strength and high modulus carbon fibers. Carbon 2015;93:81–7.
[6] Taylor LW, Dewey OS, Headrick RJ, Komatsu N, Peraca NM, Wehmeyer G, et al.
Improved properties, increased production, and the path to broad adoption of carbon nanotube fibers. Carbon 2021;171:689–94.
[7] Qian X, Zhi J, Chen L, Zhong J, Wang X, Zhang Y, et al. Evolution of microstructure and electrical property in the conversion of high strength carbon fiber to high modulus and ultrahigh modulus carbon fiber. Compos Appl Sci Manuf 2018;112:
111–8.
[8] Chae HG, Kumar S. Making strong fibers. Science 2008;319:908–9.
[9] Chae HG, Minus ML, Kumar S. Oriented and exfoliated single wall carbon nanotubes in polyacrylonitrile. Polymer 2006;47:3494–504.
[10] Sahin K, Fasanella NA, Chasiotis I, Lyons KM, Newcomb BA, Kamath MG, et al.
High strength micron size carbon fibers from polyacrylonitrile-carbon nanotube precursors. Carbon 2014;77:442–53.
[11] Park OK, Lee S, Joh HI, Kim JK, Kang PH, Lee JH, et al. Effect of functional groups of carbon nanotubes on the cyclization mechanism of Polyacrylonitrile (PAN).
Polymer 2012;53:2168–74.
[12] Park OK, Chae HS, Park GY, You NH, Lee S, Bang YH, et al. Effects of functional groups of carbon nanotubes on mechanical properties of carbon fibers. Compos B Eng 2015;76:159–66.
[13] Behabtu N, Young CC, Tsentalovich DE, Kleinerman O, Wang X, Ma AWK, et al.
Strong, light, multifunctional fibers of carbon nanotubes with ultrahigh conductivity. Science 2013;339:182–6.
[14] Headrick RJ, Tsentalovich DE, Berdegue J, Bengio EA, Liberman L, Kleinerman O, et al. Structure-property relations in carbon nanotube fibers by downscaling solution processing. Adv Mater 2018;30:1704482.
[15] Jeong HD, Kim SG, Choi GM, Park M, Ku BC, Lee HS. Theoretical and experimental investigation of the wet-spinning process for mechanically strong carbon nanotube fibers. Chem Eng J 2021;412:128650.
[16] Lee D, Kim SG, Hong S, Madrona C, Oh Y, Park M, et al. Ultrahigh strength, modulus, and conductivity of graphitic fibers by macromolecular coalescence. Sci Adv 2022;8:eabn0939.
[17] Kim SG, Choi GM, Jeong HD, Lee D, Kim S, Ryu KH, et al. Hierarchical structure control in solution spinning for strong and multifunctional carbon nanotube fibers.
Carbon 2022;196:59–69.
[18] Lee J, Lee DM, Kim YK, Jeong HS, Kim SM. Significantly increased solubility of carbon nanotubes in superacid by oxidation and their assembly into high- performance fibers. Small 2017;13:1701131.
[19] Headrick Robert J, Williams Steven M, Owens Crystal E, Taylor Lauren W, Dewey Oliver S, Ginestra Cedric J, liberman Lucy, ’akobi Asia Matatyaho Ya, Talmon Yeshayahu, Maruyama Benji, McKinley Gareth H, Hart A John, Pasquali Matteo. Versatile acid solvents for pristine carbon nanotube assembly. Sci Adv 2022;8:eabm3285.
[20] Davis VA, Parra-Vasquez ANG, Green MJ, Rai PK, Behabtu N, Prieto V, et al. True solutions of single-walled carbon nanotubes for assembly into macroscopic materials. Nat Nanotechnol 2009;4:830–4.
[21] Lagowski J. The chemistry of non-aqueous solvents: acid and aprotic solvents.
Academic press, London, p.139.
[22] Park J, Lee SH, Lee J, Lee DM, Yu H, Jeong HS, et al. Accurate measurement of specific tensile strength of carbon nanotube fibers with hierarchical structures by vibroscopic method. RSC Adv 2017;7:8575–80.
[23] Jang SY, Ko S, Jeon YP, Choi J, Kang N, Kim HC, et al. Evaluating the stabilization of isotropic pitch properties of carbon fibers. J Ind Eng Chem 2017;45:316–22.
[24] Wingert MC, Chen ZCY, Kwon S, Xiang J, Chen R. Ultra-sensitive thermal conductance measurement of one-dimensional nanostructures enhanced by differential bridge. Rev Sci Instrum 2012;83:024901.
[25] He B, Wang B, Wang Z, Qi S, Tian G, Wu D. Mechanical properties of hybrid composites reinforced by carbon fiber and high-strength and high-modulus polyimide fiber. Polymer 2020;204:122830.
[26] Zheng L, Zhang K, Liu L, Xu F. Biomimetic architectured Kevlar/polyimide composites with ultra-light superior anti-compressive and flame-retardant properties. Compos B Eng 2022;230:109485.
[27] Ebagninin KW, Benchabane A, Bekkour K. Rheological characterization of poly (ethylene oxide) solutions of different molecular weights. J Colloid Interface Sci 2009;336:360.
[28] Dubey M, Nema SK. Alloys of polyimides with varying degree of crosslinking through antimony salts: mechanical properties characterization. Polymer 1999;40:
5947–51.
[29] Bucossi AR, Cress CD, Schauerman CM, Rossi JE, Puchades I, Landi BJ. Enhanced electrical conductivity in extruded single-wall carbon nanotube wires from modified coagulation parameters and mechanical processing. ACS Appl Mater Interfaces 2015;7:27299–305.
[30] Young K, Blighe FM, Vilatela JJ, Windle AH, Kinloch IA, Deng L, et al. Strong dependence of mechanical properties on fiber diameter for polymer-nanotube composite fibers: differentiating defect from orientation effects. ACS Nano 2010;4:
6989–97.
[31] Ruan S, Gao P, Yu TX. Ultra-strong gel-spun UHMWPE fibers reinforced using multiwalled carbon nanotubes. Polymer 2006;47:1604–11.
[32] Meng J, Zhang Y, Song K, Minus ML. Forming crystalline polymer-nano interphase structures for high-modulus and high-tensile/strength composite fibers. Macromol Mater Eng 2014;299:144–53.
[33] Guo W, Liu C, Sun X, Yang Z, Kia HG, Peng H. Aligned carbon nanotube/polymer composite fibers with improved mechanical strength and electrical conductivity.
J Mater Chem 2012;22:903–8.
[34] Gao J, Itkis ME, Yu A, Bekarova E, Zhao B, Haddon RC. Continuous spinning of a single-walled carbon nanotube-nylon composite fiber. J Am Chem Soc 2005;127:
3847–54.
[35] Tambe PB, Bhattacharyya AR, Kamath SS, Kulkarni AR, Sreekumar TV, Srivastav A, et al. Structure-property relationship studies in amine functionalized multiwall carbon nanotubes filled polypropylene composite fiber. Polym Eng Sci 2012;52:
1183–94.
[36] Wu ML, Chen Y, Zhang L, Zhan H, Qiang L, Wang JN. High-performance carbon nanotube/polymer composite fiber from layer-by-layer deposition. ACS Appl Mater Interfaces 2016;8:8137–44.
[37] Newcomb BA, Chae HG, Gulgunje PV, Gupta K, Liu Y, Tsentalovich DE, et al. Stress transfer in polyacrylonitrile/carbon nanotube composite fibers. Polymer 2014;55:
2734–43.
[38] Davijani AAB, Chang H, Liu HC, Luo J, Kumar S. Stress transfer in nanocomposites enabled by poly(methyl methacrylate) wrapping of carbon nanotubes. Polymer 2017;130:191–8.
[39] Zhang X, Liu T, Sreekumar TV, Kumar S, Hu X, Smith K. Gel spinning of PVA/
SWNT composite fiber. Polymer 2004;45:8801–7.
[40] AnTh, nemann AF, Ruland W. Microvoids in polyacrylonitrile fibers: a small-angle X-ray scattering study. Macromolecules 2000;33:1848–52. u¨.
[41] Smarsly B, Antonietti M, Wolff T. Evaluation of the small-angle x-ray scattering of carbons using parameterization methods. J Chem Phys 2002;116:2618.
[42] Chae HG, Minus ML, Rasheed A, Kumar S. Stabilization and carbonization of gel spun polyacrylonitrile/single wall carbon nanotube composite fibers. Polymer 2007;48:3781–9.
[43] Tang D, Marchesini FH, Cardon L, D’hooge DR. State of the-art for extrudate swell of molten polymers: from fundamental understanding at molecular scale toward optimal die design at final product scale. Macromol Mater Eng 2020;305:2000340.
[44] Peng S, Shao H, Hu X. The effect of precursor properties on the strength of carbon fiber. Int J Polym Mater Polym Biomater 2004;53:601–8.
[45] Zhang MY, Niu HQ, Qi SL, Tian GF, Wang XD, Wu DZ. Structure evolutions involved in the carbonization of polyimide fibers with different chemical constitution. Mater Today Commun 2014;1:1–8.
[46] Franklin RE. Crystallite growth in graphitizing and non-graphitizing carbons. Proc.
R Soc A: Math Phys Eng Sci 1951;209:196–218.
[47] Zhou G, Liu Y, He L, Guo Q, Ye H. Microstructure difference between core and skin of T700 carbon fibers in heat-treated carbon/carbon composites. Carbon 2011;49:
2883–92.
[48] Newcomb BA. Processing, structure, and properties of carbon fibers. Compos Appl Sci Manuf 2016;91:262–82.
[49] Kim SG, Heo SJ, Kim JG, Kim SO, Lee D, Kim M, et al. Ultrastrong hybrid fibers with tunable macromolecular interfaces of graphene oxide and carbon nanotube for multifunctional applications. Adv Sci 2022:2203008.
[50] Qiu L, Zhang X, Guo Z, Li Q. Interfacial heat transport in nano-carbon assemblies.
Carbon 2021;178:391–412.
[51] Cai JY, Li Q, Easton CD, Liu C, Wilde AL, Veitch C, et al. Surface modification of carbon fibers with ammonium cerium nitrate for interfacial shear strength enhancement. Compos B Eng 2022;246:110173.
