Rheology Modulation and Melt Strength Enhancement Enable Continuous Melt Spinning of PEK-C Fibers

1. Introduction

Epoxy resin-based composites have been utilized more extensively in the aerospace industry, automotive manufacturing and other areas of engineering due to their excellent properties such as high strength, excellent thermal resistance and dimensional stability [1,2,3]. However, the highly crosslinked network structure of cured epoxy resins produces brittleness and poor fatigue resistance [4,5], which limits their application in extreme service working environment especially at high impact loads and recurring mechanical loads [6,7]. The enhancement of toughness of epoxy resin systems without decreasing their mechanical and thermal properties has therefore emerged as an important research topic [8,9,10].

Incorporating high-performance thermoplastic polymer fibers as the reinforcing phase is an effective method for enhancing the toughness of epoxy resins-based composites [11,12,13]. Among various thermoplastic polymers, phenolphthalein poly (aryl ether ketone) (PEK-C) has become of considerable interest due to its specific molecular structure, that is aromatic rings, ether linkages, and ketone groups [14,15]. PEK-C has excellent mechanical strength, thermal stability and chemical resistance [16,17]. Moreover, PEK-C is highly compatible with epoxy resins and has been widely used as a toughener in epoxy composites [18,19,20]. A thermoplastic-thermoset continuous phase structure is created when PEK-C and epoxy resin are mixed after curing. The interface between the two phases can prevent crack propagation. In addition, the thermoplastic phase absorbs energy through yielding deformation, which enhances the toughness of epoxy resin [21,22,23]. However, as the loading of PEK-C toughener increases, the solubility and fluidity of epoxy resins significantly decreased, which severely restricts the processability of epoxy resin [24,25]. The interaction and blending of PEK-C with epoxy matrices remain a focal point for toughening mechanisms [26,27,28]. In contrast, incorporating PEK-C fibers into the epoxy matrix can not only increase the mechanical strength and interlaminar toughness of the composites but also maintain structural stability during the molding and processing of the composites. Therefore, PEK-C fibers can serve as high performance reinforcement materials in advanced composite applications [29,30].

Currently, PEK-C fibers are primarily fabricated by solution spinning and electrospinning techniques. However, both methods are time-consuming and involve extensive use of volatile organic solvents in the process which pollutes the environment and demands a recovery cost of the solvents and time-consuming manufacturing process [31,32]. Both methods fail to meet the demands of industrial large-scale fabrication and applications because of the low production efficiency [33,34,35]. In contrast, melt spinning is a solvent-free, environmentally friendly and highly efficient fiber forming technology that has many benefits to mass production of the industry [36]. However, there is still a severe challenge with continuing the melt spinning of PEK-C fibers because of its molecular structure. PEK-C molecule structure consists of aromatic rings and bulky phenolphthalein side groups that result in the high molecular chain entanglements and high melt viscosity and poor flow during their melt processing. At the same time, its melt strength is insufficient to ensure stable filament formation during melt spinning, resulting in poor spinnability [37,38].

This study proposes a melt spinning strategy based on the regulation of rheological properties and enhancement of melt strength to achieve the continuous melt spinning of PEK-C fibers. A series of PEK-C samples with gradient molecular weights were synthesized by precisely adjusting the molar ratio of monomers during polymerization. The influence of molecular weight on the rheological properties of PEK-C was systematically investigated by capillary rheology tests. According to the rheological analysis, PEK-C with molecular weight of 2.01×10⁴ g mol^–1 was selected for melt spinning. The corresponding optimal spinning temperature was determined to be 370 °C. Subsequently, PEK-C powder was processed into pellets and fed into the melt spinning machine whereby continuous melt spinning PEK-C pre-oriented fibers were obtained. The results demonstrated that the PEK-C fibers had excellent mechanical properties, with a tensile strength of 0.56 cN/dtex and a breaking strength of 1.06 cN/dtex, respectively. Moreover, the spinning process parameters were optimized based on the mechanical properties and molecular chain orientation of PEK-C fibers. This work successfully achieves the continuous melt spinning of PEK-C fibers and provides an important material foundation and technical framework for the development of high-performance, environmentally friendly composites.

2. Materials and Methods

2.1. Materials

Phenolphthalein (PHT, >99.5%) and the 4,4′-difluorobenzophenone (DFK, >99.5%) were obtained from Shanghai Haohong Biomedical Technology Co., Ltd., Tetramethylene sulfone (TMS, >99.5%) was supplied by Liaoning Guanghua Chemical Co., Ltd., Dried sodium carbonate (Na₂CO₃, >99.5%) and xylene (XYL, >99.5%) were purchased from Shanghai Aladdin Chemicals (Shanghai, China)

2.2. Preparation of PEK-C

Figure 1 showed the flowchart of PEK-C polymerization process. PEK-C was synthesized via a nucleophilic substitution polycondensation reaction. For a representative DFK: PHT molar ratio of 1.044, the reaction was carried out in a three-necked flask. equipped with a mechanical stirrer, a thermometer, and a water separation apparatus. 18.15 g of anhydrous sodium carbonate, 30 g of DFK, 41.92 g of PHT, 30 mL of xylene and 220.5 g of Tetramethylene sulfone were added; the solid content of the resulting solution was 25%. With continuous flowing nitrogen and stirring, the reaction mixture was heated to 185 °C, then refluxed with water separation for 2.5 h, opened the water separation apparatus and remove the xylene. Subsequently, the temperature was maintained at 220 °C for 6.5 h to complete the reaction. The final reaction mixture was slowly poured into deionized water, yielding white, strip-shaped products. After being soaking in water for 10 h., the products were pulverized. Subsequently, the products were heated and stirred in deionized water for 1 h, then filtered; this process was repeated eight times until the filtrate was clear and colorless. Finally, the purified samples were dried in a forced-air oven at 180 °C for 4 h, yielding white polymer powder.

2.3. Melt Spinning of PEK-C

The powder of PEK-C was placed in a vacuum oven and dried at 150 °C under -0.1 MPa for 12 h. Then, a twin-screw extruder was employed for extrusion pelletizing, with the barrel temperatures set to 300 °C (Barrel 1), 350 °C (Barrels 2–3), 360 °C (Barrel 4), 365 °C (Barrel 5), and 370 °C (Barrels 6–7). The die head temperature was maintained at 370 °C. The dried powder was fed into the extruder hopper and was conveyed by the screw. Continuous cylindrical strands were extruded through the die head and transported to the pelletizer. To get yellow PEK-C pellets, the toughened strands were sliced into pellets in the shape of cylinders measuring about 3 mm in length and 2mm in diameter..

The PEK-C pellets were dried in a nitrogen-filled oven for 6 h before melt spinning process. Subsequently, the melt spinning of PEK-C was performed with a homemade melt spinning machine. A circular spinneret with a diameter of 0.50 mm and a length-to-diameter ratio of 5/0.5 was selected for the spinning assembly. Temperatures of different parts of the spinning machine were set to 358 °C (the C1 section of screw), 370 °C (the C1-C4 sections of screw), 372 °C (melt tube, spinning pack and pump). The winding drum speed was maintained at 600 m/min. By adjusting the rotation speed of the metering pump to 21 and 24 rpm, respectively, two types of pre-oriented fibers were obtained.

2.4. Characterization

The ¹H NMR spectra were acquired using a spectrometer (Bruker 600 AVANCE III, Bruker Corporation, Karlsruhe, Germany) with deuterated dimethyl sulfoxide (DMSO-d₆) as the solvent. The FT-IR spectrum was recorded using a spectrometer (Nicolet 6700, Thermo Scientific Inc., Waltham, MA, USA) over the wavenumber range of 4000–400 cm^–1; the tests were carried out in ATR mode. The molecular weight and distribution of the polymers were determined using a multi detector GPC/SEC system (1260 Infinity II, Agilent Technologies Co., Ltd., Santa Clara, CA, USA) with tetrahydrofuran (THF) as the solvent. The signals were detected by refractive index detector and evaluated using polystyrene reference calibration. The thermal weight loss behavior of the polymer was characterized by TGA (TGA 550, TA Instruments, New Castle, DE, USA). The tests were performed under a nitrogen atmosphere with a heating rate of 10 °C/min from room temperature up to 800 °C. The real-time weight loss variation with temperature was monitored to analyze the thermal decomposition temperature and high temperature resistance. The glass transition temperature (Tg) and thermal cross-linking reaction temperature were tested by DSC (Discovery DSC 250, TA Instruments, New Castle, DE, USA). DSC measurements for Tg were conducted in a nitrogen atmosphere with a heating rate of 10 °C/min. The sample weighing 5-10 mg was heated to eliminate thermal history and followed by a cooling and reheating cycle to record Tg during the second heating scan. DSC measurements for thermal cross-linking reaction were performed under a nitrogen atmosphere with a heating rate of 20 °C/min. The sample weighing 5-10 mg was heated to targeted temperature and maintained 30 min to monitor the thermal cross-link reaction. The melt flowability was measured using MFI tester (FBS-400B, Xiamen Fubusi Testing Equipment Co., Ltd., Xiamen, China) under test conditions of 360 °C with a 5 kg load. The intrinsic viscosity of the polymer was measured using an Ubbelohde viscometer tester (SBQ81834, Shanghai Xingao Instrument Co., Ltd., Shanghai, China) with a capillary inner diameter of 0.5 mm. A 0.5 g/dL polymer solution was prepared using N-methyl-2-pyrrolidone (NMP) as the solvent, and the measurement was conducted at 25 °C. Capillary rheology tests were carried out using a capillary rheometer (Capillary Rheometer RG25, Gottfert Ltd., Buchen, Germany). During the tests, a capillary die with dimension of 5/0.5 mm was employed. Rheological tests were conducted at 350, 355, 360, 365, and 370 °C, over a shear rate range of 1000–10000. The birefringence (Δn) of the fibers was measured using an SSY-KN birefringence tester (Shanghai Kailidi New Material Technology Co., Ltd., Shanghai, China). The instrument was focused until a clear image of the fiber was obtained. The CTB optical compensator was inserted into the optical path, and its compensation angle was adjusted to record angles θ₁ and θ₂, which corresponded to the complete extinction states achieved by rotating the compensator in opposite directions. Δn was subsequently calculated from the absolute value of the difference between θ₁ and θ₂ and the average fiber diameter. The mechanical properties of the fiber multifilament were tested using a YG029A fully automatic single-yarn strength tester produced by Changzhou Dahua Electronic Instrument Co., Ltd. (Changzhou, China). The test standard is based on the national standard GB/T 14344-2022 test methods for tensile Properties of Chemical Fiber Filaments [39].

3. Results and Discussion

3.1. The Polymerization Reaction of PEK-C

In this study, PEK-C was synthesized via nucleophilic substitution polycondensation, with the reaction equation shown in Figure 2. During the polymerization, excess 4, 4′-difluorobenzophenone (DFK) was used for end-capping.

Five PEK-C samples with different molecular weights were prepared by adjusting the molar ratio of DFK to phenolphthalein (PHT), as listed in Table 1. The PEK-C products were sequentially named PEK-C-1, PEK-C-2, PEK-C-3, PEK-C-4, and PEK-C-5 in ascending order of the molar ratio.

The molecular structure of PEK-C was characterized by proton nuclear magnetic resonance (¹H NMR) spectroscopy, and the results are showed in Figure 3a. The high symmetry of the PEK-C molecular structure and the overlap of aromatic proton peaks complicate the assignment of the relatively simple spectral lines. On this basis, the assignment of each proton signal peak was performed mainly based on the spectral line intensity in this study. The sum of the integrated proton numbers of all characteristic peaks in the full spectral range was calculated to be 20, which is completely consistent with the total number of hydrogen atoms in a single repeating unit of PEK-C. The ¹H NMR results confirm the successful synthesis of the PEK-C polymer.

The molecular structure of the polymer was further characterized by Fourier transform infrared (FT-IR) spectroscopy. The FT-IR spectra of PEK-C samples with different molecular weights are shown in Figure 3b. The analysis of the FT-IR spectrum showed that the absorption peak at 1230 cm⁻¹ is attributed to the stretching vibration of Ar-O-Ar; the bands at 1495 cm⁻¹ and 1590 cm⁻¹ correspond to the skeleton vibration of the benzene ring; the peak at 1651 cm⁻¹ is ascribed to the stretching vibration of C=O in the DFK group; and the peak at 1768 cm⁻¹ originates from the stretching vibration of C=O in the PHT group. The FTIR results further verify the successful synthesis of the PEK-C polymer.

Figure 3. PEK-C polymers: (a) ¹H NMR spectrum; (b) FT-IR spectra.

Figure 3. PEK-C polymers: (a) ¹H NMR spectrum; (b) FT-IR spectra.

3.2. The Effect of PEK-C Molecular Weight on Flowability

The molecular weight and its distribution of PEK-C were characterized by gel permeation chromatography (GPC), and the results are summarized in Table 2. All PEK-C have a polydispersity index (PDI) of 1.46–1.51, indicating that it is a relatively narrow molecular weight distribution. As the molar ratio of DFK to PHT increased from 1.044 to 1.056, the number-average molecular weight (Mn) of PEK-C decreased from 2.21×10⁴ to 2.01×10⁴, and the weight-average molecular weight (Mw) decreased from 3.32×10⁴ to 3.04×10⁴. GPC results demonstrate that the molecular weight of PEK-C can be regulated by adjusting the molar ratio of DFK to PHT during polymerization. Because extreme reduction of the molecular weight would result in loss of PEK-C properties, the optimal molecular weight should be chosen according to the specific process conditions that are place in the actual melt spinning process.

The flowability of the polymers was characterized by intrinsic viscosity and melt flow index (MFI), and the results are listed in Figure 4. As the molar ratio of DFK to PHT increased from 1.044 to 1.056, the intrinsic viscosity of PEK-C decreased from 0.34 dL/g for PEK-C-1 to 0.24 dL/g for PEK-C-5. The MFI of PEK-C increases significantly with the increase of the molar ratio, from 39.45 g/10 min to 96.61 g/10 min. The presence of rigid aromatic rings and bulky phenolphthalein side groups in the PEK-C molecular chain leads to PEK-C exhibits excellent heat resistance and mechanical properties but limits its melt flowability. Therefore, the precise control of PEK-C flowability by adjusting the molar ratio of DFK to PHT during polymerization is of great significance for melt processing. The results showed the flowability of PEK-C can be precisely regulated in a wide range by adjusting the molar ratio.

The growth of polymer molecular chains relies on the equimolar reaction between the monomers DFK and PHT. When the DFK monomer is in excess, PHT is consumed completely in advance, and the excess DFK acts as an end-capping agent in the late stage of the polymerization reaction, making the molecular chains lose reactivity, thus terminating the chain growth process. With the increase of the excess degree of DFK, the end-capping effect is enhanced, the relative molecular weight of the polymer decreased, and the entanglement degree of the molecular chains is significantly reduced. According to the Mark-Houwink equation:

$$\left[\eta \right]=K{M}^a$$

(1)

In the equation, $\left[\eta \right]$is the intrinsic viscosity; $M$ is the molecular weight; within a certain range of molecular weights, $K$ and $a$ are constants. As the molecular weight decreased, the intrinsic viscosity of PEK-C decreased. Meanwhile, the lower molecular weight and reduced molecular chain entanglement degree significantly improve the mobility of the polymer molecular chains during melt processing, resulting in a substantial increase in flowability of the PEK-C melt, which is finally manifested as a significant increase in MFI and a distinct improvement in flowability.

The molar ratio of monomers provides an important experimental basis for the melt spinning processing of PEK-C. Specifically, PEK-C resin with high flowability can be prepared by increasing the molar ratio of DFK to PHT, to meet the processing requirements for melt flowability.

3.3. Thermal Properties and Thermal Stability of PEK-C

During melt processing, too low melting temperature results in inadequate flowability which is not desirable in the processing, which is unfavorable for processing; whereas too high melting temperature can also cause thermal degradation of PEK-C, which reduces their mechanical properties. Therefore, it is important to choose the right range of melt processing temperatures based on the thermal properties and flowability of PEK-C.

The differential scanning calorimeter (DSC) and thermogravimetric analysis (TGA) were used to study the thermal properties of PEK-C, and the results are shown in Figure 5. Figure 5a presents the second heating DSC curves of PEK-C with different molecular weights. Figure 5a showed that all samples had an endothermic step of glass transition (Tg) at about 220 °C and no obvious melting endothermic peak was detected indicating that PEK-C is a typical amorphous polymer without crystalline melting temperature (Tm). Figure 5b showed the Tg curves of PEK-C at different molecular weights. The results showed that all PEK-C samples have a small weight loss rate at temperature range from room temperature to 440 °C, which confirms the excellent thermal stability of PEK-C. When temperature exceeds 440 °C, all curves show a rapid weight loss step which corresponds to the thermal degradation of the PEK-C molecular chain. At 800 °C, the weight loss rates of all PEK-C samples are 43.0–44.0%, which exhibiting good thermal resistance. Figure 5c showed the TG-DTG curve of PEK-C-4. According to the TG curve, the weight loss rate of PEK-C is less than 1.0% in temperature range of 50 °C to 440 °C, indicating that good thermal stability of PEK-C. When the temperature exceeds 440 °C, the mass loss rate decreased significantly, and the weight loss rate reaches 56.6% when the temperature is further raised to 800 °C, showing good heat resistance. The DTG curve indicates that the thermal decomposition rate of PEK-C increases significantly as the temperature up to 440 °C, and the highest rate of thermal decomposition is reached at 520 °C, indicating that the molecular chain of PEK-C significantly degraded at 520 °C.

The glass transition temperature (Tg) and 5% weight loss rate (Td_5%) of PEK-C with different molecular weights are listed in Table 3. The results showed that the glass transition temperatures of all PEK-C are near 220 °C, the high Tg of PEK-C arises from the fact that the aromatic rings, carbonyl groups and PHT groups of molecules, which hinder the segmental motion of molecular chains. Moreover, the Tg of PEK-C only decreased slightly with the reduction of the polymer molecular weight. The reason is that the mobility of polymer molecular chains is mainly affected by the end groups, which account for an extremely low proportion in the long polymer chains. As an amorphous polymer, the minimum temperature of PEK-C for melt processing is set to Tg + 100 °C, which is 320 °C.

During melt processing, to avoid thermal decomposition of the polymer due to long-term residence in the processing equipment and the resulting performance degradation, the processing temperature must be much lower than Td_5% to reserve a sufficient safety margin. As shown in Table 3, with the decreased in PEK-C molecular weight, the Td_5% only shows a slight fluctuation, ranging from 484.0 °C to 493.9 °C. The reason is that most thermal decomposition reactions of polymers are initiated from the thermally unstable chain end sites, which account for an extremely low proportion in the long polymer chains, and the proportion of thermal degradation initiated by end groups is negligible. Therefore, the molecular weight of PEK-C has a negligible actual effect on Td_5%.

Table 3. Tg and Td_5% of PEK-C.

Polymer	Tg (°C)	Td5% (°C)
PEK-C-1	221.1	486.0
PEK-C-2	220.5	488.3
PEK-C-3	219.7	489.2
PEK-C-4	219.0	493.9
PEK-C-5	217.2	484.8

Table 3. Tg and Td_5% of PEK-C.

Polymer	Tg (°C)	Td5% (°C)
PEK-C-1	221.1	486.0
PEK-C-2	220.5	488.3
PEK-C-3	219.7	489.2
PEK-C-4	219.0	493.9
PEK-C-5	217.2	484.8

During melt processing, it was found that PEK-C undergoes thermal crosslinking degradation reaction at a temperature much lower than Td_5%. The molecular chains are interconnected to form a crosslinked network structure, which transforms PEK-C from thermoplastic to thermosetting. The thermal crosslinking degradation reaction will lead to a significant decrease in melt flowability, which is unfavorable for melt processing. Since the heating residence time of the polymer melting in the equipment during melt spinning processing doesn’t exceed 20 min, the thermal stability of PEK-C after keeping at different melting temperatures for 30 min was investigated to clarify the upper limit of the safe temperature for the melt spinning processing of PEK-C.

PEK-C samples were placed in a muffle furnace and held at 360, 370, 380, 390, and 400 °C for 30 min in an air atmosphere. The changes of PEK-C after heating for 30 min at different temperatures are shown in Figure 6. When the temperature is lower than 370 °C, PEK-C maintains a uniform and transparent color. When the temperature is higher than 380 °C, the degree of carbonization and blackening of PEK-C increases with the rise of temperature. In addition, the gas released by thermal decomposition is trapped inside the melt to form bubbles due to the viscosity and surface tension of the melt, and the number of bubbles increases with the increase of temperature.

Figure 7 shows the DSC curves of PEK-C during isothermal keeping for 30 min in a nitrogen atmosphere after heating to different temperatures. The results showed when the isothermal temperature is 330 °C, 340 °C, 350 °C, 360 °C, and 370 °C, the DSC curve quickly stabilizes during the isothermal process after the heat flow abrupt change from the heating stage to the isothermal stage, and no obvious characteristic endothermic or exothermic peaks appear during the entire 30 min isothermal period. The results indicate that in the temperature range of 330 °C to 370 °C, no significant chemical reactions such as thermal decomposition, crosslinking, and depolymerization occur. PEK-C maintains excellent thermal stability under long-term isothermal conditions, with no obvious deterioration of properties, which can meet the requirements of melt spinning processing for melt thermal stability. When the isothermal temperature rises to 380 °C, the DSC curve shows a significant shift to the endothermic direction during the isothermal stage, presenting a continuous endothermic effect. The total enthalpy change ΔH during the isothermal process is calculated to be 4.08 J/g by integration. As an amorphous thermoplastic resins, PEK-C hasn’t crystalline melting phase transition process. Therefore, this continuous endothermic signal does not originate from physical phase transition but corresponds to the endothermic reaction of thermal crosslinking degradation of the polymer. Therefore, the temperature during melt processing of PEK-C must not exceed 380 °C.

During melt processing, PEK-C was fed into the processing equipment and heated rapidly to a constant temperature, the residence time in the equipment is less than 20 minutes, which is highly consistent with the test procedure of rapid heating to a constant temperature followed by isothermal holding for 30 min in the DSC test. Therefore, the isothermal DSC test results are more consistent with the actual thermal history and can more accurately reflect the thermal stability of the material in the actual processing scenario.

The outcome of the above experiment proves that PEK-C has no significant thermal degradation within 30 min at 370 °C, and its thermal stability fully meets melt processing requirements. The lowest temperature for melt processing of PEK-C is determined to be 320 °C based on the previous DSC test results. The optimal melt processing temperature range of PEK-C is 320–370 °C. At such a temperature range, not only can the amorphous PEK-C melts possess adequate flow ability and formability, but it will also not experience any thermal degradation and thermal crosslinking degradation reactions during melt processing.

3.4. Rheological Properties of PEK-C

During melt spinning, the polymer melt undergoes directional flow under the action of shear force, and its rheological behavior directly determines the continuity of melt spinning, the uniformity of fiber diameter, and the final molding properties of the fiber, which is the core basis for evaluating the spinnability of materials. The melt rheological behavior of polymer melt at high shear rates in capillary rheological tests is highly consistent with that in the melt spinning process. Therefore, in this study, a capillary rheometer was used to investigate the effects of molecular weight and melt processing temperature on the shear rheological behavior of PEK-C melts. Combined with key rheological parameters including non-Newtonian index, viscous flow activation energy and structural viscosity, this work provides theoretical and experimental support for the melt spinning of PEK-C.

Figure 8a shows the curves of shear stress versus shear rate for PEK-C with different molecular weights. In the shear rate range of 1000–10000 s^–1, the shear stress of all PEK-C samples exhibited a monotonically increasing trend with the increase of shear rate, showing significant non-Newtonian flow characteristics. At the same shear rate, the shear stress decreased with the reduction of PEK-C molecular weight. This is attributed to the increase of polymer molecular weight leading to longer molecular chain length and higher entanglement degree of molecular chains, which raises the flow resistance that needs to be overcome during the melt flow, thus manifesting as higher shear stress.

To achieve continuous and stable operation of melt spinning, the melt is required to have sufficient fluidity in the shear rate range of 3000–6000 s⁻¹, and the extrusion pressure shall not exceed the safety upper limit of 15 MPa. According to the Stokes equation:

$$\Delta p\cdot \pi R^2=2\pi R\cdot L\cdot {\tau }_w$$

(2)

where Δp is the extrusion pressure (Pa), R is the radius of the capillary die (mm), L is the length of the capillary die (mm), and τ_w is the shear stress at the capillary wall (Pa).

The diameter-length ratio of the capillary die is L/R=5/0.25, its length is L=5 mm and its radius is R=0.25 mm. The calculation shows that the largest permissible shear stress of the melt corresponding to the maximum safety pressure of the equipment is 375 kPa. In the shear rate range of 3000–6000 s^–1, only sample PEK-C-5 has a shear stress of 346 kPa at 3000 s^–1, which meets the safety extrusion pressure requirement. The shear stress of other PEK-C samples all exceeds the safety upper limit of the equipment, making it impossible to achieve melt spinning. Figure 8b presents the curves of apparent shear viscosity versus shear rate for PEK-C with different molecular weights. In the shear range of 1000–10000 s⁻¹, the apparent shear viscosity of all PEK-C samples decreased with the increase of shear rate, showing typical shear thinning behavior. This is because at low shear rates, PEK-C molecular chains are in a thermodynamic equilibrium state of random entanglement, resulting in high viscosity. With the increase of shear rate, the external force from the shear flow field causes the oriented movement of molecular chains along the flow direction and the disentanglement of molecular chains, which ultimately manifests as a continuous decrease in apparent shear viscosity with the increase of shear rate. At the same shear rate, the apparent shear viscosity decreased with the reduction of molecular weight, which is consistent with the variation law of shear stress.

Melt processing temperature is the key parameter regulating the fluidity of polymer melts. Limited by the thermal crosslinking degradation reaction of PEK-C, five temperature points (350, 355, 360, 365 and 370 °C) were selected in this study to conduct capillary rheological tests on PEK-C-5, and the results are shown in Figure 8c and Figure 8d.

Figure 8c shows the curves of shear stress versus shear rate for PEK-C-5 at different melt temperatures. In the temperature range of 350–370 °C, the shear stress of PEK-C-5 increased with the increase of shear rate. At the same shear rate, the shear stress of PEK-C-5 melt decreased significantly with the increase of melt temperature. This is because the rise of temperature enhances the mobility of molecular chain segments and reduces the entanglement degree of molecular chains. Figure 8d displays the curves of shear viscosity versus shear rate for PEK-C-5 at different melt temperatures. The shear viscosity of PEK-C-5 decreased with the increase of shear rate, showing shear thinning behavior. At the same shear rate, the apparent shear viscosity of the melt decreased with the rise of melt temperature, which is consistent with the variation law of shear stress.

According to the capillary rheological test results, reducing the molecular weight of the polymer and increasing the melt processing temperature can decrease the shear viscosity of PEK-C melt, thereby meeting the process requirements of melt spinning. The melt processing temperature is limited by the thermal crosslinking degradation reaction and shall not exceed 380 °C, so the optimal melt processing temperature is determined to be 370 °C. Although reducing the molecular weight of the polymer can lower the shear viscosity of the melt, excessive reduction of molecular weight will lead to the deterioration of the mechanical properties of the polymer. PEK-C-5, with the smallest molecular weight in all the samples, just meets the extrusion pressure and fluidity requirements of melt spinning at the melt processing temperature of 370 °C and shear rate of 3000 s⁻¹.

Figure 8. Capillary rheological properties of PEK-C: (a) the curves of shear stress versus shear rate for PEK-C with different molecular weights; (b) the curves of apparent shear viscosity versus shear rate for PEK-C with different molecular weights; (c) the curves of shear stress versus shear rate for PEK-C-5 at different temperatures; (d) the curves of shear viscosity versus shear rate for PEK-C-5 at different temperatures.

Figure 8. Capillary rheological properties of PEK-C: (a) the curves of shear stress versus shear rate for PEK-C with different molecular weights; (b) the curves of apparent shear viscosity versus shear rate for PEK-C with different molecular weights; (c) the curves of shear stress versus shear rate for PEK-C-5 at different temperatures; (d) the curves of shear viscosity versus shear rate for PEK-C-5 at different temperatures.

The non-Newtonian index (n) can characterize the degree of deviation of PEK-C melt from Newtonian fluid behavior. The closer n is to 1, the closer the melt flow characteristics are to Newtonian fluid, which means it is easier to achieve stable control of the spinning process. For pseudoplastic non-Newtonian fluids, the relationship between shear stress and shear rate follows the Ostwald equation:

$${\tau }_w=K{\gamma }^n$$

(3)

where τ_w is the shear stress at the capillary wall of the melt (Pa), n is the non-Newtonian index, k is the consistency coefficient representing the consistency of the non-Newtonian fluid, and $\gamma$ is the shear rate (s⁻¹).

Figure 9a shows the curves of lgτ_w versus lg$\gamma$ for PEK-C-5 at different melt temperatures, and the slope of the fitted straight line is the non-Newtonian index. The results of PEK-C-5 at different melt temperatures are listed in Table 4. As the temperature rises from 350 °C to 370 °C, the non-Newtonian index of PEK-C-5 increases gradually from 0.548 to 0.671, indicating that increasing the temperature can significantly weaken the non-Newtonian characteristics of PEK-C melt, make its flow behavior closer to Newtonian fluid, and reduce the sensitivity of melt viscosity to shear rate. This is conducive to reducing the problems such as uneven fiber diameter and filament breakage caused by shear rate fluctuations during spinning and improving the stability of the spinning process.

The structural viscosity index ($\Delta \eta$) is an important parameter characterizing the degree of entanglement and structuring degree of molecular chains in the melt. For pseudoplastic fluids with shear thinning behavior, $\Delta \eta$ > 0. A smaller $\Delta \eta$ indicates a lower entanglement degree of molecular chains and a weaker structuring degree of the melt, thus leading to better spinnability and stability of the melt. The calculation formula of the structural viscosity index $\Delta \eta$ is as follows:

$$\Delta \eta =-\left({\displaystyle \frac{d{\gamma }^{1/2}}{d{\ln \eta }_a}}\right)\times {10}^2$$

(4)

Figure 9b shows the curves of ln${\eta }_a$ versus ${\gamma }^{1/2}$for PEK-C-5 melt at different melt temperatures. After calculations, the negative value of the slope of the straight line is $\Delta \eta$. The calculated results of the structural viscosity index of PEK-C-5 at different temperatures are listed in Table 4. According to the results, the $\Delta \eta$ of PEK-C-5 melt shows a continuous downward trend with the rise of temperature, indicating that increasing the temperature can effectively reduce the entanglement degree between molecular chains, and significantly improve the spinnability of the material. The high entanglement degree of molecular chains caused by the rigid aromatic rings in the main chain and the bulky phenolphthalein side groups of PEK-C is the core reason for the poor spinnability of PEK-C. Increasing the melt temperature to 370 °C can effectively reduce the entanglement degree of molecular chains, providing favorable conditions for continuous and stable spinning.

Figure 9. Rheological parameters of PEK-C-5: (a) the curves of lgτ_w versus lg$\gamma$ for PEK-C-5 at different melt temperatures; (b) the curves of ln${\eta }_a$ versus ${\gamma }^{1/2}$ for PEK-C-5 melt at different melt temperatures; (c) the curves of ln${\eta }_a$ versus 1000/T for PEK-C-5 at different shear rates.

Figure 9. Rheological parameters of PEK-C-5: (a) the curves of lgτ_w versus lg$\gamma$ for PEK-C-5 at different melt temperatures; (b) the curves of ln${\eta }_a$ versus ${\gamma }^{1/2}$ for PEK-C-5 melt at different melt temperatures; (c) the curves of ln${\eta }_a$ versus 1000/T for PEK-C-5 at different shear rates.

Table 4. The n and $\Delta \eta$ coefficients of PEK-C-5 at different temperatures.

Temperature (°C)	n	$\mathit{\Delta }\mathit{\eta }$
350	0.548	2.007
355	0.592	1.896
360	0.623	1.829
365	0.633	1.782
370	0.671	1.753

Table 4. The n and $\Delta \eta$ coefficients of PEK-C-5 at different temperatures.

Temperature (°C)	n	$\mathit{\Delta }\mathit{\eta }$
350	0.548	2.007
355	0.592	1.896
360	0.623	1.829
365	0.633	1.782
370	0.671	1.753

At the same shear rate, the viscous flow activation energy (Eη) is a key parameter characterizing the sensitivity of polymer melt viscosity to temperature changes. The larger the Eη, the more significant the influence of temperature fluctuations on melt viscosity, and the higher the precision requirement for temperature control during spinning. When the melt processing temperature is higher than the viscous flow temperature of the polymer (Tg + 100 °C), the relationship between the shear viscosity of the polymer and temperature follows the Arrhenius empirical formula:

$${\eta }_a=A\cdot \exp \left({\displaystyle \frac{E_{\eta }}{R\cdot T}}\right)$$

(5)

where ${\eta }_a$ is the apparent shear viscosity (Pa·s), A is the pre-exponential factor, Eη is the viscous flow activation energy (J/mol), R is the gas constant (8.314 J/(mol·K)), and T is the absolute temperature (K).

Figure 9c shows the curves of ln${\eta }_a$ versus 1000/T for PEK-C-5 at different shear rates, the Eη of PEK-C-5 at different shear rates is calculated by fitting the slope of the straight line, and the results are listed in Table 5. The results showed that PEK-C-5 has a high viscous flow activation energy at low shear rates, indicating that its viscosity is relatively sensitive to temperature changes. This is attributed to the presence of rigid aromatic rings in the main chain of PEK-C, as well as the bulky phenolphthalein side groups, which reduce the flexibility of the molecular chains and result in an obvious steric hindrance effect. With the increase of shear rate, Eη shows an overall downward trend. In the spinning interval with shear rate of 3000-6000 s⁻¹, Eη stabilizes at 55-63 kJ/mol, indicating that under high shear action, the entanglement degree of polymer molecular chains decreased, and the influence of temperature on melt viscosity is reduced, which improves the error tolerance to temperature fluctuations and is conducive to the stability of the spinning process.

Based on the above rheological result, the optimal process parameters for PEK-C melt spinning are finally determined. PEK-C-5 is selected for melt spinning because its melt has low viscosity, flow characteristics closest to Newtonian fluid and the lowest structural viscosity index. The optimal spinning temperature is 370 °C. At this temperature, PEK-C has good thermal stability and low viscosity.

3.5. Study on the Spinnability of PEK-C and Investigation of Spinning Processes

The flowchart of PEK-C fiber forming processes were shown in Figure 10. Prior to spinning, PEK-C pellets were fed into the extruder without installing the spinning pack, and the melt was extruded at a low screw speed for material testing and purging.

Table 6 summarizes the spinnability of PEK-C samples with different molecular weights. PEK-C-1 and PEK-C-2 showed no spinnability, as their excessively high melt viscosity prevented the extrudate from forming continuous filaments. PEK-C-3 and PEK-C-4 could form filamentous extrudates, but the filaments were prone to breakage with non-uniform diameters, indicating poor spinnability. PEK-C-5 produced continuous fibers upon extrusion, confirming that its molecular weight and melt viscosity were optimal for spinning.

Fibers were subsequently prepared using PEK-C-5 with the spinning pack installed. Following the experimental procedure detailed in Section 2.3, uncollected as-spun fibers and pre-oriented fibers were obtained.

The drawing process was shown in Figure 10, the hot drawing rollers effect fiber drawing via the linear speed differential between the preceding and subsequent roller pairs, with the linear speed of the second pair of hot drawing rollers being higher than that of the first pair. The linear speed of the hot setting rollers is slightly lower than that of the drawing rollers, permitting a minor shrinkage of 1–5% in the fibers. This process eliminates the residual internal stress induced by hot drawing and markedly enhances the dimensional stability and thermal resistance of the fibers.

Figure 10. Flowchart of PEK-C fiber forming processes.

Figure 10. Flowchart of PEK-C fiber forming processes.

The study investigates the mechanical properties and molecular chain orientation of PEK-C at different drawing ratios, revealing the regulation mechanisms of fiber structure and properties during the spinning-drawing process. PEK-C pre-oriented fibers were subjected to drawing at different draw ratios, yielding three sets of PEK-C fibers with draw ratios of 1.5, 2.0 and 2.5, designated as DR-1.5, DR-2.0 and DR-2.5.

Their mechanical properties were characterized via tensile testing, and the tensile properties of PEK-C fibers are summarized in Table 7. Specifically, the pre-oriented PEK-C fibers exhibited a tensile strength of merely 0.367 cN/dtex and a breaking strength of 0.490 cN/dtex. With increasing draw ratio, the tensile strength of the drawn PEK-C fibers increased from 0.499 cN/dtex to 0.560 cN/dtex, while the breaking strength rose from 0.643 cN/dtex to 1.059 cN/dtex. Both properties showed a continuous upward trend with increasing draw ratio. This demonstrates that hot drawing treatment exerts a significant optimization effect on the mechanical properties of PEK-C fibers.

Birefringence measurements were performed on PEK-C fibers with different draw ratios, and the results are presented in Table 8. The pre-oriented PEK-C fibers showed no effectively detectable birefringence signal because pre-oriented fibers are undrawn fibers formed by natural cooling and solidification of melt filaments, in which the internal molecular chains are in a randomly coiled and entangled state with no significant preferential orientation along the fiber axis. The birefringence Δn of PEK-C fibers increased from 0.07 to 0.09 with increasing draw ratio. This structural feature is consistent with the low mechanical strength of pre-oriented fibers: randomly oriented molecular chains cannot achieve efficient axial stress transfer, resulting in poor load-bearing capacity of the fibers. The increase in Δn of PEK-C fibers with draw ratio indicates that the tensile stress along the fiber axis can induce disentanglement and axial orientation of PEK-C molecular chains, and the degree of orientation continuously improves with increasing draw ratio.

Combined with the mechanical property results, the increasing trend of Δn of PEK-C fibers is completely consistent with the rising trends of tensile strength and breaking strength with draw ratio and corresponds to the decreasing trend of elongation at break. During drawing, the high axial orientation of molecular chains enables efficient stress transfer along the rigid molecular backbones under external force, fully exploiting the intrinsic mechanical properties of PEK-C polymer and ultimately leading to a significant improvement in fiber strength. Meanwhile, the conformational space available for movement of highly oriented molecular chains is greatly restricted, reducing the plastic deformation capacity of the fibers. Consequently, the elongation at break decreased as the degree of orientation increased. The birefringence test results of this study reveal the core mechanism by which hot drawing regulates the mechanical properties of PEK-C fibers at the aggregated state structure level.

4. Conclusions

In summary, continuous melt spinning of PEK-C fibers was successfully achieved through rheology modulation and melt strength enhancement. By adjusting the monomer ratio, a series of PEK-C resins with gradient molecular weights were synthesized, enabling precise control of melt viscosity while maintaining the thermal stability and mechanical property. The PEK-C with a molecular weight of 2.01×10⁴ g mol^–1 exhibited suitable melt fluidity and processability, and the optimal spinning temperature was determined to be 370 °C. Further compounding of as-polymerized PEK-C powder into pellets allows stable and continuous melt spinning for fiber fabrication. The obtained PEK-C fibers exhibited tensile strength and breaking strength of 0.56 and 1.06 cN dtex^–1, respectively. Moreover, the relationship between spinning parameters, mechanical properties, and molecular chain orientation was further clarified, providing guidance for optimizing the melt spinning process of PEK-C fibers. This work establishes a feasible strategy for the continuous and green manufacturing of PEK-C fibers, which is expected to resolve the environmental contamination and low productivity issues associated with traditional solution spinning and electrospinning.