Facile Fabrication of PA66 / GO / MWNTs-COOH Nanocomposites and Their Fibers

: Good dispersion and interfacial compatibility are the key issues to realize the full potential of the physical–mechanical properties of nanocarbon-materials reinforced composites. Styrene–maleic-anhydride-copolymer (SMA)-treated graphene oxide (GO), carboxylated multiwalled carbon nanotubes (MWNTs-COOH), and solid-state shear milling (S3M) were applied to further improve the physical–mechanical properties of the nanocomposite ﬁbers. The results show that a mixture of GO / MWNTs-COOH exhibits good dispersion and interfacial compatibility in polyamide-66 (PA66) matrix. Consequently, the physical–mechanical properties of the ﬁbers, which were spun from the nanocomposite of GO / MWNTs-COOH treated using SMA and S3M methods, show a signiﬁcant enhancement compared to the untreated ﬁbers as well as better crystallization and thermal properties. In particular, the tensile strength of the PA66 / GO / MWNTs-COOH nanocomposite ﬁbers with a loading of 0.3 wt % GO / MWNTs-COOH reaches a maximum (979 MPa), which is the highest among all of the reported literature values. Moreover, the ﬁbers were fabricated by a facile process with e ﬃ ciency, holding great potential for industrial applications.

Composite materials with graphene and carbon nanotube additives have long been considered as exciting prospects among nanotechnology applications. In order to improve physical-mechanical and electrical properties, many researchers have paid more attention to the graphene-and carbon-nanotube-reinforced polymers [15][16][17][18]. The improvement of physical-mechanical properties of the nanocomposite materials is expected to fill the gap in the performance and price between carbon-fiberand glass-fiber-reinforced resin composites [19]. However, graphene oxide (GO) and multiwalled carbon nanotubes (MWNTs) are difficult to disperse uniformly in polymer matrix due to their large specific surface area and ratio of length/diameter, high surface energy, and strong π-π conjugate forces, which make them stack and entangle to an extreme degree. Additionally, the weak surface binding energy

Fabrication of the SMA-Treated Mixture of GO/MWNTs-COOH
GO (2 g) and MWNTs-COOH (3 g) powder were added to SMA aqueous solution (200 g) to obtain a uniformly dispersed suspension, sonicated at 50 • C for 4 h. The suspension was separated by filtration. The black solid was dried by an air-circulating oven at 25 • C for 24 h. The black powder, with a mass ratio of 2:3 between GO and MWNTs-COOH, was obtained by grinding.

Fabrication of the PA66/GO/MWNTs-COOH Nanocomposite Chips
PA66 chips were first dried by using an oven at 130 • C for 10 h. Then, PA66 chips and 5 wt % GO/MWNTS-COOH mixtures were fed into self-designed S3M equipment with a rotation rate of 525 rpm. It took 5 min to complete a whole cycle of milling 5 kg of the mixture of PA66 and GO/MWNT-COOH. A fine compounding powder of PA66 and 5 wt % GO/MWNTs-COOH can be obtained by milling five cycles. By diluting the compounding powder with pure PA66 through a melting process (two-rotor continuous extrusion processing at 280 • C), the PA66/GO/MWNTs-COOH nanocomposites (with various loading of GO/MWNTs-COOH: 0.1, 0.2, 0.3, 0.4, and 0.5 wt %) were fabricated. The same extrusion process was used in the conventional blending nanocomposites, in which the same GO/MWNTs-COOH loadings with pure PA66 chips were mixed without the S3M process.

Fabrication of the PA66/GO/MWNTs-COOH Nanocomposite Fibers
All the chips were dried and tumbled in a drum dryer to remove moisture at 130 • C for 24 h under vacuum. By doing so, the moisture on the surface of the chip was less than 50 ppm, which was measured by a Karl Fischer moisture meter. Figure 1 shows the fabrication of the nanocomposite and the melt spinning process. Chips were melted at 290 • C and then extruded with a multifilament spinneret (24 holes) with a diameter of 0.3 mm. First, the primary fiber was stretched five times using the rapid drawing process, and four godet rollers were used (GR 1: 85 • C, 70 m/min; GR 2: 170 • C, 280 m/min; GR 3: 210 • C, 350 m/min; GR 4: 80 • C, 350 m/min). Then, the obtained fiber was further stretched 1.2 times to its final length by the slow drawing process, and three sets of godet rollers were used (GR 5: 70 • C, 20 m/min; GR 6: 150 • C, 23 m/min; GR 7: 200 • C, 24 m/min). The draw ratio between the first godet roller (GR 1) and the final godet roller (GR 7) was 6, with the final yarn specification being 125 dtex/24 filaments.

Fabrication of the PA66/GO/MWNTs-COOH Nanocomposite Fibers
All the chips were dried and tumbled in a drum dryer to remove moisture at 130 °C for 24 h under vacuum. By doing so, the moisture on the surface of the chip was less than 50 ppm, which was measured by a Karl Fischer moisture meter. Figure 1 shows the fabrication of the nanocomposite and the melt spinning process. Chips were melted at 290 °C and then extruded with a multifilament spinneret (24 holes) with a diameter of 0.3 mm. First, the primary fiber was stretched five times using the rapid drawing process, and four godet rollers were used (GR 1: 85 °C, 70 m/min; GR 2: 170 °C, 280 m/min; GR 3: 210 °C, 350 m/min; GR 4: 80 °C, 350 m/min). Then, the obtained fiber was further stretched 1.2 times to its final length by the slow drawing process, and three sets of godet rollers were used (GR 5: 70 °C, 20 m/min; GR 6: 150 °C, 23 m/min; GR 7: 200 °C, 24 m/min). The draw ratio between the first godet roller (GR 1) and the final godet roller (GR 7) was 6, with the final yarn specification being 125 dtex/24 filaments.

Characterization
Fourier-transform infrared (FTIR) spectra of the samples were identified by a spectrometer (Bruker TERSOR37, Karlsruhe, Germany) in the range of 4000-400 cm −1 . The nanocomposite fibers were characterized by a field-emission scanning electron microscope (FE-SEM, Hitachi S-4800, Tokyo, Japan). The microstructures of the nanocomposite fibers were observed using a transmission electron microscope (TEM, Hitachi H-7650, Tokyo, Japan) operating at an accelerating voltage of 100 kV. The thermal properties of 5-10 mg samples were identified using a differential scanning calorimeter (DSC, Netzsch 204 F1, Bavaria, Germany). DSC scans were performed from 30 to 300 • C with a rate of ±10 • C/min in a heating-cooling-heating cycle in N 2 atmosphere. A STA409PC system from NETZSCH (Bavaria, Germany) was used to test thermogravimetric analysis (TGA) and derivative thermogravimetry (DTG) under a flowing nitrogen atmosphere, where samples (5-10 mg) were heated from 40 to 800 • C at a heating rate of 10 • C/min. A dynamic mechanical analyzer (NETZSCH DMA 242E, Bavaria, Germany) was used to measure the glass-transition temperature (T g ) of the nanocomposite fibers at a vibration frequency of 1 Hz and scanning temperature range from 30 • C to 120 • C at a heating rate of 3 • C/min. A diffractometer (Rigaku D/MAX-gA, Tokyo, Japan) with filtered Cu Kα radiation (λ = 0.15406 nm) was used for X-ray diffraction (XRD) in a range from 3 • to 40 • (2θ) and a scan rate of 8 • /min at room temperature. An optical microscope (OM, OLYMPUS-BX51, Tokyo, Japan) was used to characterize the dispersibility and secondary agglomeration of GO/MWNTs-COOH in the PA66/GO/MWNTs-COOH nanocomposite fiber with various GO/MWNTs-COOH loadings. A fiber tensile tester (ZS408, Huaxin, China) was used to test the physical-mechanical properties of the PA66/GO/MWNTs-COOH nanocomposite fibers with a test length of 200 mm at a rate of 200 mm/min for 10 times.
The molecular weight of the pure PA66 and the nanocomposites chips were calculated from the Mark-Houwink Equation (1): where M is the average molecular weight and [η] is the intrinsic viscosity (K = 3.93 × 10 −2 mL/g, α = 0.78).
[η] was measured at a concentration of 0.01 g/mL in H 2 SO 4 with a flexible system (AVS 370, Julabo, Bavaria, Germany) at 20 • C. The relative viscosity (η r ) of samples was measured three times.

FTIR Analysis
The FTIR spectra of the GO, MWNTs-COOH, and GO/MWNT-COOH powders are presented in

Viscosity Analysis
In order to assess the impact of the loaded GO/MWNTs-COOH on the molecular weight of PA66, the molecular weight of PA66 and GO/MWNTs-COOH/PA66 was determined through relative viscosity measurements, as shown in Table 1. It can be seen that the relative viscosity of PA66 with the GO/MWNTs-COOH loading increases first and then decreases. As the incorporation of GO/MWNTs-COOH increased, the relative viscosity of GO/MWNTs-COOH/PA66 increased from 2.64 to 2.86 when 0.3 wt % GO/MWNTs-COOH was loaded. This phenomenon can be described as follows: When the loading of GO/MWNTs-COOH is low, it can be uniformly dispersed in the PA66 matrix, which restrains the movement of PA66 chains and causes the fluidity to deteriorate, thereby increasing the viscosity. When the load is further increased, GO and MWNTs-COOH agglomerate, and the hindrance to the PA66 molecular chain is reduced-the fluidity is improved and the viscosity is lowered.

Morphology and Structure of GO/MWNTs-COOH and Nanocomposite Fibers
TEM micrographs of GO, MWNTs-COOH, and their mixtures are shown in Figure 3. Compared with pure GO and MWNTs-COOH, the mixture shows MWNTs-COOH inserting into the layers of GO, which is obtained using the SMA-treatment, thus reducing the stacking of GO and the agglomeration of MWNTs-COOH. Many folds can be observed on the surface of GO. These folds provide a better interface with the matrix, providing better stress transfer in the matrix. This result is

Viscosity Analysis
In order to assess the impact of the loaded GO/MWNTs-COOH on the molecular weight of PA66, the molecular weight of PA66 and GO/MWNTs-COOH/PA66 was determined through relative viscosity measurements, as shown in Table 1. It can be seen that the relative viscosity of PA66 with the GO/MWNTs-COOH loading increases first and then decreases. As the incorporation of GO/MWNTs-COOH increased, the relative viscosity of GO/MWNTs-COOH/PA66 increased from 2.64 to 2.86 when 0.3 wt % GO/MWNTs-COOH was loaded. This phenomenon can be described as follows: When the loading of GO/MWNTs-COOH is low, it can be uniformly dispersed in the PA66 matrix, which restrains the movement of PA66 chains and causes the fluidity to deteriorate, thereby increasing the viscosity. When the load is further increased, GO and MWNTs-COOH agglomerate, and the hindrance to the PA66 molecular chain is reduced-the fluidity is improved and the viscosity is lowered.

Morphology and Structure of GO/MWNTs-COOH and Nanocomposite Fibers
TEM micrographs of GO, MWNTs-COOH, and their mixtures are shown in Figure 3. Compared with pure GO and MWNTs-COOH, the mixture shows MWNTs-COOH inserting into the layers of GO, which is obtained using the SMA-treatment, thus reducing the stacking of GO and the agglomeration of MWNTs-COOH. Many folds can be observed on the surface of GO. These folds provide a better interface with the matrix, providing better stress transfer in the matrix. This result is advantageous for enhancing the binding between GO/MWNTs-COOH and PA66 and for improving the mechanical properties of the fiber.  Solving the agglomeration problem of GO/MWNTs-COOH and ensuring the homogeneous dispersion of GO/MWNTs-COOH along the nanocomposite fibers' axes in polymer matrices are the key to improve the mechanical performance of the nanocomposite fibers. Figure 4a-d are SEM images of the surface and the fracture surface of pure PA66 and nanocomposite fibers. The GO and MWNTs-COOH are evenly dispersed, both GO and MWNTs-COOH are apt to adhere on PA66 chains for effectively transferring the stress on the matrix to GO and MWNTs-COOH, leading to GO/MWNTs-COOH resisting to be pulled out from the matrix. Thereby, it is hoped that the physical-mechanical properties of the nanocomposite fibers are effectively improved. The morphology of the nanocomposite fibers was observed by OM, in which the dispersion of dark spots is the GO/MWNTs-COOH, consistent with the SEM microimages. Moreover, the GO/MWNTs-COOH without agglomeration dispersed homogeneously in polymer matrices along the nanocomposite fiber's axis can be observed in Figure 4f. The orientation of GO/MWNTs-COOH along the drawing direction is crucial for the mechanical performance of the nanocomposites [33]. Therefore, the SEM and OM micrograph results indicate that the nanocomposite fibers with 0.3 wt % of GO/MWNTs-COOH have a better mechanical performance. Solving the agglomeration problem of GO/MWNTs-COOH and ensuring the homogeneous dispersion of GO/MWNTs-COOH along the nanocomposite fibers' axes in polymer matrices are the key to improve the mechanical performance of the nanocomposite fibers. Figure 4a-d are SEM images of the surface and the fracture surface of pure PA66 and nanocomposite fibers. The GO and MWNTs-COOH are evenly dispersed, both GO and MWNTs-COOH are apt to adhere on PA66 chains for effectively transferring the stress on the matrix to GO and MWNTs-COOH, leading to GO/MWNTs-COOH resisting to be pulled out from the matrix. Thereby, it is hoped that the physical-mechanical properties of the nanocomposite fibers are effectively improved. The morphology of the nanocomposite fibers was observed by OM, in which the dispersion of dark spots is the GO/MWNTs-COOH, consistent with the SEM microimages. Moreover, the GO/MWNTs-COOH without agglomeration dispersed homogeneously in polymer matrices along the nanocomposite fiber's axis can be observed in Figure 4f. The orientation of GO/MWNTs-COOH along the drawing direction is crucial for the mechanical performance of the nanocomposites [33]. Therefore, the SEM and OM micrograph results indicate that the nanocomposite fibers with 0.3 wt % of GO/MWNTs-COOH have a better mechanical performance.   Figure 5 and Table 2 show the physical-mechanical properties of the nanocomposite fibers. Due to the even distribution of GO/MWNTs-COOH within the polymer matrix, and due to interfacial interaction [34], it is obvious that the mechanical behavior of the nanocomposite fiber with GO/MWNTs-COOH in a polymer matrix is enhanced. The tensile strength of 0.3 wt % GO/MWNTs-COOH nanocomposite fiber with S3M process is 979 MPa, which is an improvement of 32% compared to a pure PA66 fiber. The stress fiber tends to increase with lower loading, and elongation remains constant. When GO/MWNTs-COOH loading further increases, from 0.3 wt % to 0.5 wt %, the tensile strength decreases from 979 to 766 MPa. The physical-mechanical properties of the nanocomposite fibers without the S3M decrease after an initial increase. However, the performance of the fiber without the S3M process is much less than that of the fiber with the S3M process. It can be proved that the S3M process makes GO/MWNTs-COOH disperse evenly in PA66. With increased GO/MWNTs-COOH loadings in PA66, a reassembly occurs due to van der Waals forces interacting  Figure 5 and Table 2 show the physical-mechanical properties of the nanocomposite fibers. Due to the even distribution of GO/MWNTs-COOH within the polymer matrix, and due to interfacial interaction [34], it is obvious that the mechanical behavior of the nanocomposite fiber with GO/MWNTs-COOH in a polymer matrix is enhanced. The tensile strength of 0.3 wt % GO/MWNTs-COOH nanocomposite fiber with S3M process is 979 MPa, which is an improvement of 32% compared to a pure PA66 fiber. The stress fiber tends to increase with lower loading, and elongation remains constant. When GO/MWNTs-COOH loading further increases, from 0.3 wt % to 0.5 wt %, the tensile strength decreases from 979 to 766 MPa. The physical-mechanical properties of the nanocomposite fibers without the S3M decrease after an initial increase. However, the performance of the fiber without the S3M process is much less than that of the fiber with the S3M process. It can be proved that the S3M process makes GO/MWNTs-COOH disperse evenly in PA66. With increased GO/MWNTs-COOH loadings in PA66, a reassembly occurs due to van der Waals forces interacting between nanosheets. Due to the agglomerated GO/MWNTs-COOH, the physical-mechanical properties of the nanocomposite fibers decrease, resulting in stress concentration. With increasing GO/MWNTs-COOH loadings, the Young's modulus of the nanocomposite fibers gradually increases compared to the pure fiber. Until 0.5 wt % GO/MWNTs-COOH loading, Young's modulus is improved by about 41% compared to the pure fiber. The reason may be attributed to the interaction between GO/MWNTs-COOH and the PA66, which limits the motion of the PA66 chains. With increasing GO/MWNTs-COOH loading, the elongation of the nanocomposite fibers is almost constant. Compared to the S3M-untreated samples, the S3M-treated samples have better physical-mechanical properties at the same GO/MWNTs-COOH loading. This indicates that the S3M process can improve the polymer performance. Since the crystallization and orientation of the fiber were further improved by the two-step drawing process, better physical-mechanical properties can be obtained. According to PA66 fibers reinforced by graphene, carbon nanotubes, and other additives [4,21,35,36], our work has the highest physical-mechanical properties among all of the reported values. between nanosheets. Due to the agglomerated GO/MWNTs-COOH, the physical-mechanical properties of the nanocomposite fibers decrease, resulting in stress concentration. With increasing GO/MWNTs-COOH loadings, the Young's modulus of the nanocomposite fibers gradually increases compared to the pure fiber. Until 0.5 wt % GO/MWNTs-COOH loading, Young's modulus is improved by about 41% compared to the pure fiber. The reason may be attributed to the interaction between GO/MWNTs-COOH and the PA66, which limits the motion of the PA66 chains. With increasing GO/MWNTs-COOH loading, the elongation of the nanocomposite fibers is almost constant. Compared to the S3M-untreated samples, the S3M-treated samples have better physicalmechanical properties at the same GO/MWNTs-COOH loading. This indicates that the S3M process can improve the polymer performance. Since the crystallization and orientation of the fiber were further improved by the two-step drawing process, better physical-mechanical properties can be obtained. According to PA66 fibers reinforced by graphene, carbon nanotubes, and other additives [4,21,35,36], our work has the highest physical-mechanical properties among all of the reported values.

Thermal and Crystallization Behavior
TGA and DTG tests were conducted for the pure and the nanocomposite fibers. The mass loss, due to degradation volatilization, was monitored as a function of temperature, as shown in Figure 6, corresponding to a 5% mass loss (T 5 wt % ). The temperatures at the maximum rate of mass loss are  Table 3. With GO/MWNTs-COOH additions increasing, the thermal stability of the nanocomposite fibers was improved first and then reduced. Here, T max is the temperature at the maximum mass loss rate regarded as the decomposition temperature. T max of the nanocomposite fiber with 0.3 wt % GO/MWNTs-COOH loading is 445.2 • C, which is higher than that of the pure fiber (439.4 • C). The T max improvement of the nanocomposite fiber is ascribed to the compatibility of GO/MWNTs-COOH and the PA66 matrix, which makes heat transfer fast and improves the heat resistance of the nanocomposite fiber. TGA and DTG tests were conducted for the pure and the nanocomposite fibers. The mass loss, due to degradation volatilization, was monitored as a function of temperature, as shown in Figure 6, corresponding to a 5% mass loss (T5 wt %). The temperatures at the maximum rate of mass loss are summarized in Table 3. With GO/MWNTs-COOH additions increasing, the thermal stability of the nanocomposite fibers was improved first and then reduced. Here, Tmax is the temperature at the maximum mass loss rate regarded as the decomposition temperature. Tmax of the nanocomposite fiber with 0.3 wt % GO/MWNTs-COOH loading is 445.2 °C, which is higher than that of the pure fiber (439.4 °C). The Tmax improvement of the nanocomposite fiber is ascribed to the compatibility of GO/MWNTs-COOH and the PA66 matrix, which makes heat transfer fast and improves the heat resistance of the nanocomposite fiber.  The DSC results of the pure and the nanocomposite fibers with various GO/MWNTs-COOH loadings are shown in Figure 7. Compared with the pure PA66 fibers, Tm of the nanocomposite fibers increases first and then decreases. This is explained by the GO/MWNTs-COOH and PA66 that disperse evenly and form strong interfacial interactions, resulting in a more rapid heat transfer by SMA treatment and the S3M process. Upon further increasing the GO/MWNTs-COOH loading, reassembling causes Tm to decrease.  Table 3. TGA parameters of the pure and the nanocomposite fibers. The DSC results of the pure and the nanocomposite fibers with various GO/MWNTs-COOH loadings are shown in Figure 7. Compared with the pure PA66 fibers, T m of the nanocomposite fibers increases first and then decreases. This is explained by the GO/MWNTs-COOH and PA66 that disperse evenly and form strong interfacial interactions, resulting in a more rapid heat transfer by SMA treatment and the S3M process. Upon further increasing the GO/MWNTs-COOH loading, reassembling causes T m to decrease.

Loadings (wt %)
The crystallization temperature (T c ) can be determined by analyzing cooling scans. Compared with the pure PA66 fibers, T c of the nanocomposite fibers with 0.3 wt % GO/MWNTs-COOH loading increases by 7.2 • C due to the heterogeneous nucleating action of GO/MWNTs-COOH. The GO/MWNTs-COOH loading can improve the crystallization rate by reducing the formation power of a new crystal plane. The degree of crystallinity (X c ) values, calculated with various GO/MWNTs-COOH loadings from the normalized heat of fusion of the DSC, are listed in Table 4. The maximum crystallinity appears at 0.3 wt % GO/MWNTs-COOH loading, which is higher than that of the pure PA66 fiber. When further increasing the GO/MWNTs-COOH loading, GO/MWNTs-COOH reassembling restricts the movement of PA66 molecular chains, resulting in incomplete crystallization and decreased crystallinity. By comparing Table 2 with Table 4, the strength and crystallinity present a positive correlation, i.e., the higher the crystallinity, the higher the strength.  Table 4. The maximum crystallinity appears at 0.3 wt % GO/MWNTs-COOH loading, which is higher than that of the pure PA66 fiber. When further increasing the GO/MWNTs-COOH loading, GO/MWNTs-COOH reassembling restricts the movement of PA66 molecular chains, resulting in incomplete crystallization and decreased crystallinity. By comparing Table 2 with Table 4, the strength and crystallinity present a positive correlation, i.e., the higher the crystallinity, the higher the strength. The loss tangent (tan δ) of the pure PA66 fiber and the PA66/GO/MWNTs-COOH nanocomposite fibers with various GO/MWNTs-COOH loadings are shown in Figure 8. Due to the loss factor being positively correlated with the amorphous area of the fiber, as GO/MWNTs-COOH loading increases, from 0 to 0.3 wt %, the loss factor reduces from 0.104 to 0.058. It can also be seen that both the crystalline region and crystallinity of the fiber increase. Compared with Tg of the pure PA66 fiber (55.3 °C), Tg of 0.3 wt % GO/MWNTs-COOH nanocomposite fibers rises, reflected from the loss tangent peaks of the nanocomposite fibers shifting toward high temperature, increasing to 62.7 °C. Tg of nanocomposite fibers is enhanced due to GO/MWNTs-COOH with strong affinity dispersed evenly within the PA66 matrix, limiting polymer segmental motion and reducing the free volume.  The loss tangent (tan δ) of the pure PA66 fiber and the PA66/GO/MWNTs-COOH nanocomposite fibers with various GO/MWNTs-COOH loadings are shown in Figure 8. Due to the loss factor being positively correlated with the amorphous area of the fiber, as GO/MWNTs-COOH loading increases, from 0 to 0.3 wt %, the loss factor reduces from 0.104 to 0.058. It can also be seen that both the crystalline region and crystallinity of the fiber increase. Compared with T g of the pure PA66 fiber (55.3 • C), T g of 0.3 wt % GO/MWNTs-COOH nanocomposite fibers rises, reflected from the loss tangent peaks of the nanocomposite fibers shifting toward high temperature, increasing to 62.7 • C. T g of nanocomposite fibers is enhanced due to GO/MWNTs-COOH with strong affinity dispersed evenly within the PA66 matrix, limiting polymer segmental motion and reducing the free volume.    Figure 9 shows XRD patterns of pure PA66 and the PA66/GO/MWNTs-COOH nanocomposite fibers. PA66 is a polycrystalline polymer with α and β crystal shapes. The α-PA66 (100) (including β-PA66), (010), and (110) crystal surface diffraction peaks can be demarcated at diffraction angles 20.4 • and 23.4 • . The relative peak height of the (100), (010), and (110) peaks reduce with increased GO/MWNTs-COOH loading. This phenomenon is because crystals tend to grow along the (010) and (110) crystal surfaces when GO/MWNTs-COOH are added to PA66. According to Scherrer's equation, the grain size of the pure PA66 and the composite fiber can be obtained, as shown in Table 5. It shows that the grain size of the (100) crystal surface of the composite fiber is less than that of PA66, and those of the (010) and (110) crystal surfaces of the composite fiber are more than that of PA66. This explains why the degree of crystallization is increased, as it is caused by the action of heterogeneous nucleation of GO/MWNTs-COOH in the PA66 crystallization process, facilitating crystal growth along the (010) and (110) crystal surfaces.  Figure 9 shows XRD patterns of pure PA66 and the PA66/GO/MWNTs-COOH nanocomposite fibers. PA66 is a polycrystalline polymer with α and β crystal shapes. The α-PA66 (100) (including β-PA66), (010), and (110) crystal surface diffraction peaks can be demarcated at diffraction angles 20.4° and 23.4°. The relative peak height of the (100), (010), and (110) peaks reduce with increased GO/MWNTs-COOH loading. This phenomenon is because crystals tend to grow along the (010) and (110) crystal surfaces when GO/MWNTs-COOH are added to PA66. According to Scherrer's equation, the grain size of the pure PA66 and the composite fiber can be obtained, as shown in Table  5. It shows that the grain size of the (100) crystal surface of the composite fiber is less than that of PA66, and those of the (010) and (110) crystal surfaces of the composite fiber are more than that of PA66. This explains why the degree of crystallization is increased, as it is caused by the action of heterogeneous nucleation of GO/MWNTs-COOH in the PA66 crystallization process, facilitating crystal growth along the (010) and (110) crystal surfaces.

Conclusions
The physical, mechanical, and thermal properties of melt-spun PA66/GO/MWNTs-COOH nanocomposite fibers fabricated by using styrene-maleic anhydride copolymer treatment and solid-state shear milling were improved. MWNTs-COOH integrated into the layers of GO and dispersed uniformly in the PA66 matrix, and interfacial adhesions between GO/MWNTs-COOH and the PA66 was improved under the 3D forces exerted by the S3M. The mechanical properties of the S3M-treated nanocomposite fibers are better than that of the pure, non-S3M-treated fibers. In particular, the tensile strength and Young's modulus of the S3M-treated nanocomposite fibers with a 0.3 wt % GO/MWNTs-COOH loading reach up to 979 MPa and 6.37 GPa, respectively. In addition, the effect of S3M process on the crystallization and thermal properties of nanocomposite fibers is remarkable, leading to an increase of the thermal stability of the S3M-treated fibers with 0.3 wt % GO/MWNTs-COOH loading of 5.7 • C.
The utilization of SMA-treated GO/MWNTs-COOH, the S3M process, melt spinning, and the two-step drawing process all provide a facile and cost-effective method of the direct use of GO/MWNTs-COOH for the large-scale fabrication of nanocomposite fibers with excellent thermal properties and enhanced mechanical properties, which can be widely used at the industrial scale.