Method for producing a polymer fiber and polymer fiber produced thereby

Abstract

This disclosure generally relates to a method for producing polymer fibers, typically polyacrylonitrile-based fibers, the characteristics of which are controlled by certain parameters of the method, such as the amount of stretching by the nozzle used, the amount of wet stretching, and the amount of hot stretching. This disclosure also relates to a method for producing carbon fibers from such polymer fibers.

Description

[0001] Cross-reference to related applications

[0002] This application claims priority to U.S. Provisional Application No. 63 / 157,111, filed March 5, 2021, the entire contents of which are incorporated herein by reference. Technical Field

[0003] This disclosure generally relates to a method for producing polymer fibers, typically polyacrylonitrile-based fibers, the characteristics of which are controlled by certain parameters of the method, such as the amount of stretching by the nozzle used, the amount of wet stretching, and the amount of hot stretching. This disclosure also relates to a method for producing carbon fibers from such polymer fibers. Background Technology

[0004] Carbon fiber, especially polyacrylonitrile (PAN)-based carbon fiber, is in increasing demand in many market sectors due to its high performance-to-weight ratio. 1 However, its market reach and wider industrial acceptance are limited by its relatively high cost. 2,3 To reduce costs, carbon fiber manufacturers are seeking larger tow sizes due to increased capacity, higher production volumes, and lower capital costs per pound of fiber processed. 4-7 Although large tows (>24,000 filaments) offer significant cost advantages, they are hampered by processing challenges, making wet spinning the preferred manufacturing method. 8,9 Large tows are plagued by additional downstream challenges, including a higher percentage of filament breakage during hot drawing. 8 Steam stretching or conventional winding mechanisms cannot be applied. 9,10 The high exothermicity during thermal oxidation stabilization (TOS) and the variability in carbon fiber yield. 11 Despite these additional challenges in processing large tows, the characteristic targets and performance indicators are not lower than their small tow counterparts.

[0005] In order to achieve the target properties of carbon fiber, attention has shifted to the development of precursors and fiber spinning in the early stages to establish the structure and morphology of carbon fiber. 12-15 The precursor structure is affected by chemical composition, 16 , 17 Fiber fineness and crystallinity 18 Fiber structure and arrangement 19-21 The effect of spinning speed. 20 Structural defects (such as cavities, cracks, and imperfections) imposed by spinning methods may affect fiber shrinkage during TOS and reduce the strength of the resulting carbon fibers. 22-24Therefore, elucidating the key features of wet spinning for developing PAN precursor structures can improve the mechanical properties of carbon fibers.

[0006] The morphology and structure of PAN have been clearly defined in the literature. 13 However, the researchers were initially puzzled by the different experimental observations under different methodological histories. 25-28 Bashir et al. conducted a series of concise studies and clarified that the structure of random PAN can be divided into unoriented and oriented states. 29-32 In the unoriented state, PAN exhibits a two-phase morphology represented by amorphous domains and ordered intermediate domains, as observed by two thermal transitions in differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA). 30,32 Low-temperature DMA transition ((β) c The α peak (around 100℃) is attributed to intermediate phase domains, while the high-temperature transformation (around 130-160℃) is attributed to amorphous domains. As PAN fibers become oriented and chain stacking occurs during drawing, the intensity of the α peak decreases until it disappears. At this point, the oriented PAN becomes a "laterally ordered single phase". 26,31,33,34

[0007] In addition to thermal transitions and chain stacking, the chain conformation also exhibits a transformation as the PAN transitions between unoriented and oriented states. Sawai et al. and others have demonstrated that unoriented random PANs employ a larger ratio of irregular helical sequences, and that oriented chains rearrange into a planar zigzag conformation during drawing. 35-37 Shen et al. announced a syndiotactic system with the highest degree of planar zigzag arrangement, possessing the lowest energy, the highest stiffness, and the most favorable chain conformation for intermolecular packing and TOS. 38

[0008] Although precursor structures in both unoriented and oriented states have been well-defined, the evolution of the microstructure of PAN-based precursors is rarely studied, or the pathway directly connecting unoriented to oriented states in fiber spinning is rarely investigated. This is because simulating commercial spinning processes and studying the effects of process variations on fiber structure is capital-intensive and difficult. Instead, much research has focused on the early stages of fiber spinning, where coagulation parameters can have a direct and profound impact on precursor structure. 39-44 In particular, much attention has been paid to the dosing extrusion rate and nozzle stretching, or stretching outside the coagulation between the spinneret and the first guide plate in wet spinning. 45-49

[0009] Besides nozzle stretching, wet spinning includes other stretching stages that can be used interchangeably to achieve the target filament fineness before TOS and carbonization. Stretching can be performed immediately after solidification in a gel bath or some non-solvent medium. 39,50,51 Conventional dry stretching, 37,52,53 Alternatively, it can be achieved through a hot drawing process using hot liquid or a steam box. 51 Although each pulling stage reduces the fiber diameter, the pulling mechanism can differ significantly due to the structural properties of the gel state and the thermally pulled state. Edrington's paper confirms the overall pulling limit of the combination of pulling stages, but does not delve into separating nozzle pulling and gel pulling from thermal pulling to determine where the greatest impact on crystal alignment and orientation occurs. 51

[0010] Therefore, there is an urgent need to achieve higher consistency in orientation states through optimal spinning processes that result in more stable crystalline regions and favorable precursor structures for carbon fiber conversion. This paper describes a method for producing polymer fibers suitable for carbon fiber formation. The method of this invention utilizes insights derived from studies of the microstructure of PAN-based polymer fibers via wide-angle X-ray scattering (WAXS) and an understanding of thermal transition and stabilization kinetics in DMA via DSC. Summary of the Invention

[0011] Advantageously, it has been unexpectedly discovered that a given stretching stage is irreversible, and that the stretching profile can play a significant role in the precursor structure. The method of this invention utilizes insights derived from studies of the microstructure of polymer fibers (such as PAN-based fibers) via wide-angle X-ray scattering (WAXS) and an understanding of thermal transitions and stabilization kinetics in DMA via DSC.

[0012] In a first aspect, this disclosure relates to a method for producing polymer fibers, the method comprising:

[0013] a) The polymer solution is spun into a coagulation bath, where a nozzle is applied to stretch the fibers to form coagulated fibers;

[0014] b) subjecting the solidified fibers obtained in step (a) to wet stretching to form a first drawn fiber; and

[0015] c) The first drawn fiber obtained in step (b) is subjected to thermal stretching to form a polymer fiber;

[0016] The amount of nozzle stretching, wet stretching, and hot stretching are effective in achieving the following characteristics:

[0017] The resulting polymer fibers have a Herman orientation factor of at least 0.60, typically at least 0.65, and more typically at least 0.67.

[0018] The resulting polymer fiber has a crystallite thickness that is at least 3 nm thicker than that of the first drawn fiber, typically at least 3.5 nm thicker, and more typically at least 4 nm thicker.

[0019] In a second aspect, this disclosure relates to a method for producing carbon fibers, the method comprising oxidizing the polymer fibers described herein or polymer fibers produced according to the methods described herein to form stabilized carbon fiber precursor fibers, and then carbonizing the stabilized carbon fiber precursor fibers to produce carbon fibers. Attached Figure Description

[0020] Figure 1 The structural evolution observed throughout the methods described herein, such as by WAXS analysis of the fibers, is illustrated using specific tensile curves. Detailed Implementation

[0021] As used herein, unless otherwise stated, the terms “a / an” or “the” mean “a / an or more / multiple” or “at least one / a” and are used interchangeably.

[0022] As used in this article, the term “and / or” in phrases in the form of “A and / or B” means A alone, B alone, or A and B together.

[0023] As used herein, the term "comprises" includes "consistsessentially of" and "consists of". The term "comprising" includes "consisting essentially of" and "consisting of". "Comprising", synonymous with "including", "containing", or "characterized by", is intended to be inclusive or open-ended and does not exclude additional, unlisted elements or steps. The transitional phrase "consisting essentially of" includes specific materials or steps that do not materially affect the essential characteristics or function of the described composition, process, method, or article. The transitional phrase "composed of" does not include any unspecified elements, steps, or components.

[0024] Unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this specification pertains.

[0025] As used herein, and unless otherwise specified, the terms “about” or “approximately” mean an acceptable error for a particular value as determined by one of ordinary skill in the art, which depends in part on how the value is measured or determined. In some embodiments, the terms “about” or “approximately” mean within 1, 2, 3, or 4 standard deviations. In some embodiments, the terms “about” or “approximately” mean within 50%, 20%, 15%, 10%, 9%, 8%, 7%, 6%, 5%, 4%, 3%, 2%, 1%, 0.5%, or 0.05% of a given value or range.

[0026] Furthermore, it should be understood that any numerical ranges listed herein are intended to include all subranges included therein. For example, the range "1 to 10" is intended to include and encompass all subranges between the listed minimum value of 1 and the listed maximum value of 10; that is, a minimum value equal to or greater than 1 and a maximum value equal to or less than 10. Because the disclosed numerical ranges are continuous, they include every value between the minimum and maximum values. Unless otherwise clearly indicated, the various numerical ranges specified in this application are approximate values.

[0027] Throughout this disclosure, different publications may be incorporated by reference. Unless otherwise specified, if the meaning of any language in such publications incorporated by reference conflicts with the meaning of the language in this disclosure, the meaning of the language in this disclosure shall prevail.

[0028] A first aspect of this disclosure relates to a method for producing polymer fibers, the method comprising:

[0029] a) The polymer solution is spun into a coagulation bath, where a nozzle is applied to stretch the fibers to form coagulated fibers;

[0030] b) subjecting the solidified fibers obtained in step (a) to wet stretching to form a first drawn fiber; and

[0031] c) The first drawn fiber obtained in step (b) is subjected to thermal stretching to form a polymer fiber;

[0032] The amount of nozzle stretching, wet stretching, and hot stretching are effective in achieving the following characteristics:

[0033] The resulting polymer fibers have a Herman orientation factor of at least 0.60, typically at least 0.65, and more typically at least 0.67.

[0034] The resulting polymer fiber has a crystallite thickness that is at least 3 nm thicker than that of the first drawn fiber, typically at least 3.5 nm thicker, and more typically at least 4 nm thicker.

[0035] In step a) of this method, a polymer solution (or “spinning solution”) is spun into a coagulation bath to form coagulated fibers. Typically, the polymer solution is a homogeneous solution containing a polyacrylonitrile-based polymer and a solvent. Therefore, in the examples, the resulting polymer fibers are polyacrylonitrile-based polymer fibers.

[0036] Polyacrylonitrile-based polymers can be any polymer comprising repeating units derived from acrylonitrile. Suitable polyacrylonitrile-based polymers can be homopolymers composed of repeating units derived from acrylonitrile or copolymers comprising repeating units derived from acrylonitrile and one or more comonomers. Such polymers can be obtained from commercially available sources or prepared according to methods known to those skilled in the art. For example, the polymers can be prepared by any polymerization method, including but not limited to solution polymerization, dispersion polymerization, precipitation polymerization, suspension polymerization, emulsion polymerization, and variations thereof.

[0037] Polyacrylonitrile-based polymers comprise repeating units derived from acrylonitrile and at least one comonomer selected from the group consisting of: vinyl-based acids, vinyl-based esters, vinyl amides, vinyl halides, ammonium salts of vinyl compounds, sodium salts of sulfonic acids, and mixtures thereof.

[0038] In the embodiments, the polyacrylonitrile-based polymer comprises repeating units derived from acrylonitrile and at least one comonomer selected from the group consisting of: methacrylic acid (MAA), acrylic acid (AA), itaconic acid (ITA), methacrylate (MA), ethyl acrylate (EA), butyl acrylate (BA), methyl methacrylate (MMA), ethyl methacrylate (EMA), propyl methacrylate, butyl methacrylate, β-hydroxyethyl methacrylate, dimethylaminoethyl methacrylate, 2-ethylhexyl acrylate, isopropyl acetate, vinyl acetate (VA), vinyl propionate, vinylimidazolium (VIM), acrylamide (AAm), diacetone acrylamide (DAAm), allyl chloride, vinyl bromide, vinyl chloride, vinylidene chloride, sodium vinyl sulfonate, sodium p-styrene sulfonate (SSS), sodium methyl allyl sulfonate (SMS), sodium 2-acrylamido-2-methylpropane sulfonate (SAMPS), and mixtures thereof.

[0039] The comonomer ratio (the amount of one or more comonomers relative to the amount of acrylonitrile) is not particularly limited. However, suitable comonomer ratios are 0 to 20%, typically 1% to 5%, and more typically 1% to 3%.

[0040] The molecular weight of the polyacrylonitrile-based polymer suitable for use according to the method can be in the range of 60 to 500 kg / mol, typically 90 to 250 kg / mol, and more typically 115 to 180 kg / mol.

[0041] Suitable solvents for polymers can be selected from the group consisting of: dimethyl sulfoxide (DMSO), dimethylformamide (DMF), dimethylacetamide (DMAc), ethylene carbonate (EC), N-methyl-2-pyrrolidone (NMP), zinc chloride (ZnCl2) / water, sodium thiocyanate (NaSCN) / water, and mixtures thereof, typically selected from the group consisting of: dimethyl sulfoxide (DMSO), dimethylformamide (DMF), dimethylacetamide (DMAc), ethylene carbonate (EC), and N-methyl-2-pyrrolidone (NMP).

[0042] The polymer solutions used are typically free of gelling and / or agglomerated polymers. The presence of gelling and / or agglomerated polymers can be determined using any method known to those skilled in the art. For example, the Hegman gauge can be used to determine the presence of gelling and / or agglomerated polymers. The polymer solutions used are generally stable and do not gel over time.

[0043] The concentration of the polymer in the polymer solution is suitably at least 10 wt% by weight, typically from about 16 wt% to about 28 wt%, based on the total weight of the solution.

[0044] After removing air bubbles by vacuum, the polymer solution can be subjected to conventional wet spinning and / or air-gap spinning. In wet spinning, the dopant is filtered and extruded through orifices in a spinneret (typically made of metal) into a liquid coagulation bath for the polymer to form filaments. The spinneret orifices determine the desired fiber count (e.g., 3,000 orifices for 3K carbon fibers). In air-gap spinning, a vertical air gap of 1 to 50 mm, typically 2 to 10 mm, is provided between the spinneret and the coagulation bath. In this embodiment, the polymer solution is spun by wet spinning.

[0045] The coagulation bath used in this method is a mixture of solvent and non-solvent. Water or alcohol is typically used as the non-solvent. Suitable solvents include those described herein. In examples, dimethyl sulfoxide, dimethylformamide, dimethylacetamide, or mixtures thereof are used as solvents. In another example, dimethyl sulfoxide is used as the solvent. The ratio of solvent to non-solvent and the bath temperature are not particularly limited and can be adjusted according to known methods to achieve the desired curing rate of the newly extruded filaments during coagulation. However, the coagulation bath typically contains 40 wt% to 85 wt% of one or more solvents, with the balance being a non-solvent. In examples, the coagulation bath contains a mixture of DMSO and water.

[0046] In step a), nozzle stretching is used. As used herein, the amount of nozzle stretching applied in the coagulation step is the ratio of the first roll take-up speed to the dopant extrusion speed, as will be understood by one of ordinary skill in the art. 46,48 Original liquid extrusion speed V jet The following equation 1 calculates the volumetric flow rate, typically provided by a metering pump, where N is the number of orifices in the spinneret, and D is the diameter of each orifice. Therefore, those skilled in the art will recognize that V... jet This can be adjusted by choosing appropriate values ​​for Q, N, and D.

[0047]

[0048] In step b) of the method, the solidified fibers obtained in step (a) are subjected to wet stretching to form a first drawn fiber. As used herein, wet stretching is synonymous with "wet drawing" or "first draw". st "Draw" is synonymous with "FD" and these terms are used interchangeably. During wet drawing, the coagulated fiber is immersed in one or more baths and conveyed through one or more baths, typically by the action of rollers or drawing rollers. As used herein, the amount of wet drawing is defined as the ratio of the speed of the drawing rollers between coagulation and the first drawing. The one or more baths used for wet drawing can be maintained at a temperature between 40°C and 100°C.

[0049] In step c) of the method, the first drawn fiber obtained in step (b) is subjected to thermal stretching to form a polymer fiber. As used herein, thermal stretching is synonymous with "thermal drawing" or "second draw". nd"Draw" is synonymous with "HD" and these terms are used interchangeably. During hot drawing, the first drawn fiber is conveyed through a heat source without immersion in a liquid, typically by the action of rollers or drawing rollers. For example, the heat source may be steam or a tubular furnace through which the first drawn fiber obtained in step (b) is conveyed. As used herein, the amount of hot drawing is defined as the ratio between the drawing roller speed after FD and the final winding speed.

[0050] The method disclosed herein can be performed continuously or in batches. As used herein, a "continuous" method means that in this method, fibers are conveyed through one or more processing steps, one single work unit at a time, without any interruption in time, material, or sequence. This contrasts with a batch method, which should be understood as a method comprising a series of one or more steps performed in a defined sequence, and whereby a limited quantity of material is processed or produced at the end of that sequence, and the sequence must be repeated to process or produce another batch of material. In the embodiments, the method is performed continuously.

[0051] The polymer fibers produced by the methods described herein, typically polyacrylonitrile-based fibers, can be used as precursor fibers for the production of carbon fibers, so-called white fibers (“WF”). Therefore, in the embodiments, the polymer fibers produced are carbon fiber precursor fibers.

[0052] In the method described herein, the amount of nozzle stretching, wet stretching, and hot stretching are effective in achieving the following characteristics: the Herman orientation factor of the resulting polymer fiber is at least 0.60, typically at least 0.65, more typically at least 0.67, and the crystallite thickness of the resulting polymer fiber is at least 3 nm, typically at least 3.5 nm, more typically at least 4 nm greater than the crystallite thickness of the first drawn fiber.

[0053] It has been found that solidification affects the initial structure and indicates orientation and structural evolution. Hot drawing or high-temperature drawing, followed by wet drawing, is beneficial for increasing the size of intermediate phase domains, and structural order directly affects the mechanical properties of TOS and downstream process parameters. Therefore, adjusting the amount of nozzle drawing, wet drawing, and hot drawing enables the achievement of certain properties.

[0054] The parameters were measured using wide-angle X-ray scattering (WAXS) and then used to determine various fiber properties throughout the method, including Herman orientation factor, crystallite thickness, crystallinity, and D-spacing.

[0055] WAXS testing is performed according to methods known to those skilled in the art. In a suitable instance, the Xenocs Xeuss 2.0 system is used to perform WAXS in transmission mode. The light source is... The wavelength λ of the GeniX3DCu ULD was 8 keV. A single fiber bundle (1,000 filaments) was aligned and secured along the aperture card. The aperture card with the aligned fiber bundles was then transferred to the WAXS sample holder, which was positioned 101.17 mm from the 2D detector (Dectris Pilatus 200K). The exposure time of the precursor fibers was 600 s. Data processing was performed using Foxtrot software provided by Xenocs to obtain the results relative to 2θ and azimuth angle. The integrated diffraction intensity. Structural parameters can be determined by peak fitting of Foxtrot data using the Lorentzian function with Origin 2017 software, showing two crystalline peaks located at approximately 2θ1≈16.9° and 2θ2≈29.3°, and concentrated at approximately 2θ... a Amorphous peak ≈25°. 54 Crystallinity χ is determined using Equation 2. c A θ1 and A θ2 It refers to the areas of the two crystallization peaks, 2θ1 and 2θ2, and A. θa These represent the areas of the amorphous peaks. The D-spacing is determined by Bragg's Law (Equation 3) and the crystallite thickness, L... c The shape factor K is determined by the Scherrer equation (Equation 4), where K is 0.89 and β is the full width at half maximum (FWHM) of the corresponding crystallization peak (in radians). The azimuth scan curve is obtained by azimuth integration of the diffraction intensity I in the (100) diffraction plane (2θ1 ≈ 16.9°). 55,56 Herman orientation factor (f c The answer is determined by equations 5 and 6. 39

[0056]

[0057]

[0058]

[0059]

[0060]

[0061] Herman orientation factor, crystallite thickness, crystallinity, and D-spacing can be determined at any point in the methods described herein. For example, fiber samples can be taken after each step of the method and measured by WAXS to determine the Herman orientation factor, crystallite thickness, crystallinity, and D-spacing of the fibers formed in each step.

[0062] In the examples, the Herman orientation factor of the resulting polymer fibers is at least 0.60, typically at least 0.65, and more typically at least 0.67.

[0063] In another embodiment, the Herman orientation factor of the coagulated fiber is at least 0.35, typically at least 0.40, and more typically at least 0.42.

[0064] In another embodiment, the Herman orientation factor of the resulting polymer fiber is at least 0.08 greater than that of the first drawn fiber, typically at least 0.1.

[0065] In the embodiments, the resulting polymer fiber has a crystallite thickness that is at least 3 nm greater than that of the first drawn fiber, typically at least 3.5 nm, and more typically at least 4 nm.

[0066] In the embodiments, the crystallinity of the solidified fiber is no more than 8%, typically no more than 7%, and more typically no more than 6% greater than the crystallinity of the first drawn fiber.

[0067] Throughout this method, the fibers obtained in each step can be characterized by their cross-sectional diameter. Any suitable method for measuring fiber diameter can be used. In a suitable example, scanning electron microscopy (SEM) is used. In imaging preparation, the polymer fibers are freeze-dried on a Labconco freeze dryer using a tert-butanol / water cosolvent system. The filament cross-section is produced by immersing the fiber bundle in water and then immersing it in liquid nitrogen before breakage. The filament ends are mounted on a 15 mm aluminum SEM stub using a carbon ribbon and coated with 3 nm platinum sputtering to reduce charging. The cross-section can be imaged using a Hitachi S-4800 at 2 kV and different magnifications. The average filament diameter can be measured using QUARTZ PCI image processing software. At least 10 measurements are performed for each sample.

[0068] In the embodiments, the average diameter of the coagulated fibers is at least 40 μm, typically at least 45 μm, more typically at least 50 μm, and even more typically at least 55 μm.

[0069] In another embodiment, the average diameter of the first drawn fiber is at least 15 μm, typically at least 20 μm, and typically at least 22 μm.

[0070] The fibers obtained in each step of this method can also be obtained through the mesophase glass transition temperature (β). c The activation energy ΔH of structural relaxation under the condition is used to characterize it.

[0071] The mesophase glass transition temperature (β) was measured using dynamic mechanical analysis (DMA) and methods known to those skilled in the art. c The activation energy ΔH of structural relaxation under [condition] was determined. Suitablely, DMA was performed on a TA Instruments Discovery HybridSeries HR-2 rheometer. The precursor fiber bundle was tested using a rectangular tension clamp geometry in tension mode. A 5-minute temperature immersion and an axial force of 1 ± 0.1 N were applied to condition the sample. Loss modulus, storage modulus, and tanδ variable were plotted against temperature at a rate of 1 °C / min over a temperature range of 45 °C to 165 °C at varying frequencies of 0.5, 1, 5, 10, and 15 Hz. The axial strain was set to 0.1%. The mesophase glass transition temperature (β [value]) was determined. c The activation energy ΔH of structural relaxation under ) is taken as the peak in tanδ and calculated by Equation 7 through the Arrhenius relation with frequency f and the general gas constant R. 33,35,57

[0072]

[0073] In the examples, the resulting polymer fibers have β c The activation energy for structural relaxation is less than 700 kJ / mol, typically less than 650 kJ / mol, and more typically less than 600 kJ / mol.

[0074] In another embodiment, the activation energy of the structural relaxation of the resulting polymer fibers βc ranges from 500 to 600 kJ / mol, typically from 530 to 570 kJ / mol.

[0075] The fibers obtained in each step of this method can be characterized by the cyclization activation energy. The cyclization activation energy can be measured using any method known to those skilled in the art. A suitable method is differential scanning calorimetry (DSC). For example, DSC can be performed on a TA instrument Q2000 equipped with a Universal Analysis 2000. DSC is calibrated according to ASTM E967-03 and ASTM E968-02 for temperature and heat flux, using indium metal as a certified reference material (melting temperature: 156.60 °C ± 0.03 °C, enthalpy of fusion: 28.70 J / g ± 0.09 J / g), respectively. The standard aluminum DSC pan of the TA instrument with a cap can be used with a nitrogen or air flow rate of 55 mL / min. The DSC is equilibrated at 35 °C for 2 min, and then ramped from 35 °C to 450 °C using ramp rates of 2, 5, 10, 20, and 30 °C / min.

[0076] The cyclization activation energy was determined by the Flynn-Wall-Ozawa (FWO) method (Equation 8), where E a It is the cyclization activation energy, R is the universal gas constant, and T is the cyclization activation energy. pk ω is the peak heat release temperature (in Kelvin), and ω is the temperature ramp rate (in Kelvin). 58,59

[0077]

[0078] In the embodiments, the cyclization activation energy of the resulting polymer fiber is at least 7 kJ / mol greater than that of the first drawn fiber, typically at least 11 kJ / mol, and more typically at least 13 kJ / mol.

[0079] DSC can also be used to determine the peak temperature of the cyclization exothermic reaction of the fibers obtained in each step of the method. The cyclization reaction is typically exothermic and can be viewed as a peak in a DSC scan, where heat flux (in W / g) is plotted as a function of temperature. In the examples, the peak temperature of the cyclization exothermic reaction of the resulting polymer fibers is at least 3°C ​​higher than the peak temperature of the cyclization exothermic reaction of the solidified fibers and / or the first drawn fibers.

[0080] Polymer fibers formed by the methods of the present invention described herein, typically carbon fiber precursor fibers or white fibers, possess mechanical properties such as toughness, elongation, and Young's modulus, due to unexpected observations of certain microstructures caused by the stretching profiles (i.e., the amount of nozzle stretching, wet stretching, and hot stretching).

[0081] The linear mass density, or tow fineness, of the resulting polymer fibers can be calculated by averaging two yield measurements based on the weight of a 3-meter sample portion. Linear mass density can be expressed in terms of tow (a bundle of filaments) or on a per-filament basis. In this paper, linear mass density is expressed on a per-filament basis. Therefore, the linear mass density used herein is the mass (in grams) per 9000 meters of filament, or denier / filament. The linear mass density of the resulting polymer fibers ranges from 0.7 to 1.2, typically from 0.85 to 1.0 denier / filament.

[0082] Mechanical properties can be tested using known methods. For example, MTS Criterion C43 and Testworks 4 software can be used. Fiber samples were loaded with a gauge length of 7.875 inches, a preload force of 0.1 lbf, and a clamping pressure of 60 psi. Each sample was tested a total of 3 times at a crosshead speed of 2 in / min.

[0083] In the examples, the resulting polymer fibers have a toughness of at least 4 g / d, typically at least 5 g / d, and more typically at least 6 g / d.

[0084] In the examples, the resulting polymer fibers have a Young's modulus of at least 95 g / d, and more typically at least 100 g / d.

[0085] In another embodiment, the resulting polymer fibers have a Young's modulus of 95 to 130 g / d, typically from 100 to 130 g / d, and more typically from 115 to 125 g / d.

[0086] In a second aspect, this disclosure relates to a method for producing carbon fibers, the method comprising oxidizing the polymer fibers described herein or polymer fibers produced according to the methods described herein to form stabilized carbon fiber precursor fibers, and then carbonizing the stabilized carbon fiber precursor fibers to produce carbon fibers.

[0087] The polymer fibers formed according to the methods described herein can be oxidized to form stabilized carbon fiber precursor fibers, and the stabilized carbon fiber precursor fibers can then be carbonized to produce carbon fibers.

[0088] During the oxidation stage, polymer fibers are fed under tension through one or more dedicated ovens, each with a temperature ranging from 150°C to 300°C, typically from 200°C to 280°C, wherein heated air is fed into each oven. The oxidation process binds oxygen molecules from the air to the fibers and causes the polymer chains to begin cross-linking, thereby increasing the fiber density to a suitable level. Typically, PAN fibers oxidized in this way have an infusible ladder-like aromatic molecular structure and are ready for carbonization treatment.

[0089] Carbonization leads to the crystallization of carbon molecules and thus produces finished carbon fibers with a carbon content greater than 90%. The carbonization of oxidized or stabilized carbon fiber precursor fibers occurs in an inert (oxygen-free) atmosphere, typically a nitrogen atmosphere, within one or more specially designed furnaces. The oxidized carbon fiber precursor fibers are then passed through one or more ovens, each heated to a temperature ranging from 300°C to 1650°C.

[0090] Surface treatment may involve traction carbonized fibers through an electrolytic bath containing an electrolyte such as ammonium bicarbonate or sodium hypochlorite. The chemicals in the electrolytic bath etch or roughen the fiber surface, thereby increasing the surface area available for interfacial fiber / matrix bonding and adding reactive chemical groups that can be used to form composite materials.

[0091] Next, the carbon fibers can be subjected to sizing, in which a sizing coating (e.g., an epoxy-based coating) is applied to the fibers. Sizing can be performed by passing the fibers through a sizing bath containing a liquid coating material. Sizing protects the carbon fibers during processing and finishing into intermediate forms such as dry fabrics and prepregs. Sizing also holds the filaments together in individual bundles to reduce fuzz, improve processability, and increase the interfacial shear strength between the fibers and the matrix resin used to manufacture composites. After sizing, the coated carbon fibers are dried and then wound onto a winding tube. The resulting carbon fibers are suitable for use in the production of composite materials.

[0092] The following non-limiting examples further illustrate the methods disclosed herein and the polymer fibers produced therefrom.

[0093] Example

[0094] Example 1. Preparation of polymer fibers

[0095] Polymer dope was prepared by mixing a PAN-based polymer in DMSO at a target concentration of 18.5 wt.% solids in a 15-gallon Myers mixer. The polymer dope was heated to 65°C, degassed under vacuum, and pressurized with nitrogen in a dope storage tank. The polymer dope was then fed through a pressure-regulated and heated metering pump and filtered before wet spinning. The dope flow rate was controlled at a constant rate of 40 mL / min and extruded through a 1,000-hole spinneret (orifice diameter = 50 μm) into a heated coagulation bath containing a mixture of DMSO and water.

[0096] The polymer composition, bath concentration, temperature, and other spinning parameters remained constant, with the only variable being the ratio of stretching taken during the nozzle stretching and second stretching. Three variants were investigated, with the stretching amount proportional to the amount taken between the nozzle stretching and the second stretching, as outlined in Table 1 below. The variants are designated as “1,” “2,” or “3,” and the steps from which fiber samples were taken are designated as “Coag” (after coagulation), “FD” (after the first stretching), and “WF” (after the second stretching). The amount of nozzle stretching applied during coagulation was controlled by the ratio X of the first roll take-up speed to the dope extrusion speed. 46,48 Original liquid extrusion speed V jet Calculated from Equation 1 above. The ratio of the speed of the stretching roll between solidification (“Coag”) and the first stretch (“FD”) is the amount of wet stretching Y, and the ratio between the speed of the stretching roll after FD and the speed of the final winding machine is the amount of stretching applied in the hot stretch.

[0097] Compared to Coag-1, the nozzle stretching increased by 36% and 58% for Coag-2 and Coag-3, respectively, and the second stretching decreased by the corresponding amount. In all three variants studied, the final winding speed was equal, and therefore, the total stretching applied to each white fiber precursor was equal, X*Y*Z. For each of the three variant samples collected after coagulation and the first stretching, they were allowed to dry to a constant weight before further analysis. White fiber (WF) samples were collected on the winding machine after drying during spinning and used as is.

[0098] Table 1. Samples collected along spinning threads with different stretch ratios.

[0099]

[0100] a) The ratio of the first stretching roller to the raw material extrusion speed; b) The ratio of the second stretching roll to the first stretching roll; c) The ratio of winding machine speed to the second stretching roll; d) Total stretch: Nozzle stretch * First stretch * Second stretch

[0101] Example 2. SEM Analysis

[0102] As described above, SEM analysis was performed on the fiber samples indicated in Table 1.

[0103] It was observed that as the nozzle stretch increased from Coag-1 to Coag-2 to Coag-3, the filament cross-section changed from a kidney bean shape to an oblong shape, and finally to a more rounded shape. It is known that nozzle stretch affects the shear force in the spinneret capillary, which reduces die expansion and thus reduces solvent exchange. 46,48,60 As the filaments leave the capillary, the increased nozzle stretch also reduces the filament diameter, where the larger fibril radius has a thinner and sparser skin structure, making the reverse diffusion of water and DMSO easier. 60 Compared to Coag-2 and Coag-3, which have a more pronounced skin layer, the larger diameter of Coag-1 results in a milder skin-core transition. Compared to Coag-1 and Coag-2, for Coag-3, back diffusion is sufficiently hindered to form a robust skin and impede complete coagulation of the core structure, leading to a looser, spongy characteristic. Furthermore, the higher nozzle stretch, combined with the shorter residence time in the coagulation bath, can also contribute to incomplete fiber coagulation.

[0104] After solidification, the fibers are subjected to the same draw ratio in the first drawing process, although the actual speed and residence time vary due to differences in nozzle stretching. In all three cases, the cross-sectional shape tends towards a circular filament, but FD-1 retains some similarities to kidney beans. Cross-sectional images were obtained by fracturing the sample in liquid nitrogen, where the fractured surface provides insight into defects. 61 FD-1 and FD-2 show radial patterns originating from the fiber center, indicating internal defects. FD-3 shows separation of the skin from the core, indicating that the breakage begins in the skin and propagates around the core of the filament rather than through the center. The filament skin also shows a distinct difference from FD-1, which has finer striations along the fiber axis, while FD-3 has more pronounced and deeper skin grooves due to its faster solidification rate. 62 The average filament diameter maintained the same proportional difference as solidification after the first drawing, which expected a constant first drawing ratio. For example, the average filament diameters of Coag-2 and FD-2 (49.6 μm and 19.1 μm, respectively) remained approximately 86% of those of Coag-1 and FD-1 (58.0 μm and 22.2 μm, respectively). A proportional adjustment was used in the second drawing to compensate for the speed difference, ensuring that the winding speed was equal for each of the three curves. The average diameters of the solidified and first-drawn fibers, as observed by SEM, are summarized in Table 2 below. Morphological testing of WF samples by SEM was difficult due to the fiber's extensibility and is therefore not included.

[0105] Table 2.

[0106]

[0107] Example 3. WAXS Analysis

[0108] To investigate the structural differences in fibers after each stretching stage, WAXS was performed on each sample. Figure 1 The structural evolution of curve 1 is shown, where the scans at each sampling location overlap. Interestingly, Coag-1 shows a strong peak at approximately 2θ1 ≈ 17°, corresponding to the region of intermediate phase orientation in the PAN. 31 This indicates that significant crystallization occurred due to solidification and nozzle stretching. Coag-1 also reveals a broad amorphous peak around 2θ. a ≈25°, and the beginning of another higher-order diffraction plane is approximately 2θ2≈29°. 54 As the fiber was further stretched during the first drawing, FD-1 showed a slight narrowing of the 2θ1 peak and a larger intensity of the 2θ2 peak compared to Coag-1. After the second drawing, WF-1 revealed a much sharper 2θ1 diffraction peak, a very pronounced 2θ2 peak, and a relatively smaller 2θ2 peak. a The intensity has decreased.

[0109] Crystallinity (χ²) under all processing conditions c ), microcrystalline thickness (L) c ) and orientation (f c The crystallinity appears to decrease during this process; however, due to the data processing techniques used, this behavior is believed to be anomalous. Interpretations of crystallinity in the literature are often controversial. 13 The proposed method used in this paper interprets the 2θ1 and 2θ2 peaks as part of ordered domains, as observed in the WAXS peak parameters and peak fitting of various fiber samples (not shown). However, 2θ2 is related to the amorphous region 2θ a The difficulty of deconvolution leads to higher apparent crystallinity during solidification, which means that samples with higher amorphous content can artificially increase the peak area of ​​the 2θ2 peak and thus increase the overall crystallinity of the sample.

[0110] Surprisingly, the prominent 2θ1 peak in the solidified sample indicates that the mesophase domains are largely established at the start of spinning. Despite the crystallinity of 2θ1 and the d-spacing (d... 2θ1 The crystallite thickness and orientation remained relatively stable throughout the stretching process of the spinning process, but significant structural changes occurred. For all three samples, the orientation gradually increased from Coag to FD to WF, with the peak values ​​of each corresponding azimuth scan narrowing. The narrowing of the 2θ1 peak indicates that the crystallite thickness increased slightly between Coag and FD, and surged from FD to WF. This suggests that hot stretching is more efficient than wet stretching when melting ordered regions together and combining amorphous and ordered domains. Liu et al. have proposed that the initial stage of stretching is achieved through the straightening and unwinding of polymer chains. 22 This can also explain L during wet stretching. c The marginal increase. The WAXS parameters of the fiber samples are summarized in Table 3 below.

[0111] Table 3. WAXS parameters of fiber samples.

[0112]

[0113] a) f c Herman orientation factor obtained by azimuth analysis of the 2θ1 peak; b) χ c Crystallinity; C) d-space of 2θ1; d) L c : 2θ1 crystallite thickness

[0114] Example 4. Molecular Dynamics Simulation

[0115] To visualize the evolving structure revealed by WAXS data, molecular dynamics simulations were performed. Here, a PAN system consisting of twenty independent chains was subjected to axial strain to simulate a stretching process.

[0116] Schrodinger Materials Science 19-3 and OPLS3-e force-field were used for all simulations. PAN chains were constructed using the “Crosslinked Polymer” module as follows: 5000 acrylonitrile monomers and 20 initiator molecules were first placed in a periodic chamber, followed by 50 ns NPT and NVT stages at 400 K to densify and relax the cells. Virtual atoms were used to represent growing radicals. Twenty reactions were allowed per step, corresponding to the number of initiators; however, the monomers had to be within a 4.0 Å distance standard from the virtual atoms to react. After each step, the system was relaxed for 50 ps. Polymerization was carried out at 400 K and continued until the monomer conversion equaled 100%. After polymerization, the system was annealed at 400 K for an additional 50 ns NPT and NVT. Tensile simulations were performed by one-dimensionally stretching the periodic chamber by 0.1% strain for 600 steps to reach a total strain of 60%. A Poisson's ratio of 0.5 was applied to control the lateral dimensions of the chamber. Stretching was performed under NVT conditions. Skeletal torsion is analyzed after stretching simulation, where torsion is defined by the 4-atom dihedral angles from the four connected skeleton atoms.

[0117] A snapshot of the system was observed through simulation, where single chains were represented by magnified atoms to track their progression. The 0% system was omitted as it is completely amorphous and not considered a good representative of real polymers. The expanded regions of the 10% and 60% strain systems reveal local ordering in the random polymer chains, where at lower strains, a large number of cis configurations act as kinks in the chains, disrupting local ordering. Observations of the simulation results highlight different numbers of skeletal dihedral angles, with two configurations being highly preferred: trans (≈180°) and cis (≈60°). As stretching proceeds, the cis conformation transforms into the trans conformation, resulting in an expanded planar zigzag arrangement and allowing adjacent chains to align and condense, thus forming larger intermediate phase domains. As stretching continues further, the cis conformation tends to zero, and the system approaches single-phase behavior.

[0118] Example 5. DMA Analysis of Fibers

[0119] The evolution of the storage modulus (E') and loss modulus (E”) of curve 1 at a frequency scan of 10 Hz was investigated using DMA analysis as described above. Due to the challenges in accurately measuring the cross-sectional area of ​​different samples, the magnitudes of the storage modulus and loss modulus should not be considered significant. Instead, tanδ, the ratio of E” to E', will be used to compare samples, as it eliminates biases related to sample size.

[0120] Coag-1 shows two peaks in the tanδ curve, with the α-transition peak at approximately 145℃ and the β-transition peak at approximately 145℃. c The transition peak is approximately 105℃. 31,33 As the fibers become oriented in FD-1, the α-transformation disappears and the β-transformation... c The magnitude of the transition begins to increase, indicating that the α peak corresponds to the amorphous region and the β peak... c This corresponds to an ordered intermediate phase region. 31 Coag-1 also showed a narrower β compared to FD-1. c The transition peak is due to the more pronounced and larger amorphous region in the spun fibers. 63 The transition to single-phase behavior is further supported as stretching through WF-1, where there is only one β phase. c The peak exists.

[0121] β c The activation energy of the transition was determined by Arrhenius fitting, using the reciprocal of the peak temperature, 1 / β. c The results are plotted as a function of the logarithm of frequency, ln f. The results are summarized in Table 4 below.

[0122] Table 4. Peak tanδ temperature at different frequencies (1℃ / min ramp) and activation energies ΔH.

[0123]

[0124]

[0125] a) β c Peak relaxation temperature, determined by a temperature ramp of 1℃ / min and a strain of 0.1% at each given frequency; b) ΔH: as in the Arrhenius relation Determined β c The activation energy of structural relaxation, where f is the frequency and R is the universal gas constant.

[0126] As the frequency increases, the transition temperature increases and the magnitude of tanδ decreases due to the shorter response time and viscoelasticity of PAN-based fibers. The slopes of stretching curve 1 follow the order: Coag-1 > FD-1 > WF-1, which corresponds to the decrease in ΔH value as the fiber is stretched.

[0127] Example 6. DSC Analysis

[0128] DSC analysis was performed on fiber samples obtained by the methods described herein. DSC results were determined for fiber evolution in nitrogen at 20 °C / min. A distinct onset peak associated with exothermic cyclization was observed at approximately 243 °C for Coag-1, which decreased in FD-1 and was absent in WF-1. This initial peak is attributed to the larger amorphous domains in Coag-1 and FD-1 compared to WF-1. WF-1 exhibited a slower onset of exothermic reaction, possibly due to a greater mixing of ordered and amorphous domains compared to Coag-1 and FD-1, which may have more discrete phase boundaries. Cyclothermic exothermic reactions in ordered domains exhibited peaks at lower temperatures and higher heat fluxes, with larger onset peaks below 250 °C and higher internal heat generation, such as approximately 296 °C for Coag-1 and FD-1 compared to approximately 299 °C for WF-1.

[0129] Determining the activation energy E of cyclization using the Ozawa method a The reciprocal of the peak exothermic temperature, 1 / T pk The logarithm lnω relative to the heating rate is plotted, which is revealed as the peak temperature increases for a faster ramp rate. 59,64 Peak temperature and calculated E a The values ​​are summarized in Table 5 below.

[0130] Table 5. Maximum DSC peak values ​​determined in nitrogen at different ramp rates and cyclization activation energies.

[0131]

[0132] a) The peak exothermic temperature of the maximum heat flux in nitrogen for DSC at a given ramp rate; b) E a Cycloning activation energy determined by the Ozawa method

[0133] The cyclization activation energy between the solidified sample and the sample after the first pulling remained almost constant, ranging from 155 kJ / mol to 158 kJ / mol. In the case of WF-1, when E... a The most significant change occurred during hot drawing when the value increased to 169.8 kJ / mol. Cycloning began to appear, attributed to structural anomalies and defects. 64-66This can explain why the WF samples have a much higher activation energy. As demonstrated by WAXS and DMA analyses, the white fibers after hot drawing contain the largest-sized ordered domains, which leave fewer structural defects for the initiation of cyclization.

[0134] Example 7. Mechanical properties of white fibers

[0135] The mechanical properties were used to compare the final fibers and are summarized in Table 6 below.

[0136] Table 6. Mechanical properties of white fibers for various draw curves.

[0137]

[0138] The linear density of each WF sample was within 1% variability, which confirmed that each given draw curve produced filaments of the same diameter (for the same given precursor composition and packing density). Interestingly, the white fiber toughness and Young's modulus increased in the order: WF-3 < WF-2 < WF-1. This behavior is consistent with the results reported by Arbab et al., 67 where the higher second draw associated with spinneret drawing is more effective in reducing porosity and increasing fiber stiffness. Additionally, the higher toughness and modulus can be expected from the higher structural order and alignment of the polymer chains observed in WF-1.

[0139] Wet spinning of PAN-based precursors was carried out on a pilot spinning line, where the draw variables were separated from the polymer composition, bath temperature, coagulation conditions, and other process parameters. Three draw distribution variants were selected, where an increase in spinneret draw was exchanged with a proportional decrease in hot drawing, resulting in a constant white fiber fineness and winder speed. Three samples were collected for analysis at each draw curve at the following locations: coagulation, first draw, and fiber winder. As determined by SEM, WAXS, DMA, and DSC, structural evolution occurred between the sample locations, but the final white fiber sample properties showed key differences, including the toughness and modulus of the white fibers.

[0140] Increasing the spinneret draw led to higher shear forces and incomplete coagulation of the fibers, which resulted in variable cross-sectional shapes and skin-core structures. In the order of the maximum orientation and microcrystalline thickness: Coag-1 > Coag-2 > Coag-3, which directly corresponded to a longer residence time and lower shear forces during coagulation. This result propagated to the samples collected after the first draw and the samples collected at the winder, where the orientation gradually increased between each draw stage. The results showed a strong correlation between orientation and the structural relaxation of the mesophase domains, which confirmed the importance of complete coagulation and the establishment of ordered domains in the primary stage of spinning. Additionally, WF-1 with the largest hot draw ratio showed the most significant L between the FD and WF stagesc The increase corresponds well to the most significant change in cyclization behavior. The cyclization activation energy shows a relationship with L... c The correlation is good, with higher ordered domains leading to higher activation energies and delayed onset temperatures. The method of the present invention described herein uses important findings: (1) solidification affects the initial structure and indicates orientation and structural evolution; (2) hot drawing, or high-temperature drawing, followed by wet drawing, is beneficial for increasing the size of mesophase domains; and (3) structural order directly affects the mechanical properties of TOS and downstream process parameters.

[0141] It will be apparent to those skilled in the art that the conditions for carrying out the methods of the present invention described herein can be optimized based on the intended application and circumstances without departing from the spirit of this disclosure.

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Claims

1. A method for producing polymer fibers, the method comprising: a) The polymer solution is spun into a coagulation bath, where a nozzle is applied to stretch the fibers to form coagulated fibers; b) subjecting the coagulated fibers obtained in step (a) to wet stretching to form a first drawn fiber, wherein the wet stretching is performed at 40°C to 100°C; and c) The first drawn fiber obtained in step (b) is subjected to thermal stretching to form the polymer fiber; The amount of nozzle stretching, wet stretching, and hot stretching are effective in achieving the following characteristics: The resulting polymer fiber has a Herman orientation factor of at least 0.60, and The resulting polymer fiber has a crystallite thickness that is at least 3 nm greater than that of the first drawn fiber. The crystallinity of the solidified fiber is no more than 8% greater than that of the first drawn fiber; and The Herman orientation factor of the solidified fiber is at least 0.

35.

2. The method according to claim 1, wherein, The Herman orientation factor of the resulting polymer fiber is at least 0.

65.

3. The method according to claim 2, wherein, The Herman orientation factor of the resulting polymer fiber is at least 0.

67.

4. The method according to claim 1, wherein, The resulting polymer fiber has a crystallite thickness that is at least 3.5 nm greater than that of the first drawn fiber.

5. The method according to claim 2, wherein, The resulting polymer fiber has a crystallite thickness that is at least 4 nm greater than that of the first drawn fiber.

6. The method according to claim 1, wherein, The crystallinity of the solidified fiber is no more than 7% greater than that of the first drawn fiber.

7. The method according to claim 6, wherein, The crystallinity of the solidified fiber is no more than 6% greater than that of the first drawn fiber.

8. The method according to any one of claims 1-7, wherein, The linear mass density of the resulting polymer fiber ranges from 0.7 to 1.2 deniers per filament.

9. The method according to claim 8, wherein, The linear mass density of the resulting polymer fiber ranges from 0.85 to 1.0 denier per filament.

10. The method according to any one of claims 1-7, wherein, The average diameter of the solidified fibers is at least 40µm.

11. The method according to claim 10, wherein, The average diameter of the solidified fibers is at least 45 µm.

12. The method according to claim 11, wherein, The average diameter of the solidified fibers is at least 50 µm.

13. The method according to claim 12, wherein, The average diameter of the solidified fibers is at least 55 µm.

14. The method according to any one of claims 1-7, wherein, The average diameter of the first drawn fiber is at least 15 µm.

15. The method according to claim 14, wherein, The average diameter of the first drawn fiber is at least 20 µm.

16. The method according to claim 15, wherein, The average diameter of the first drawn fiber is at least 22 µm.

17. The method according to any one of claims 1-7, wherein, The Herman orientation factor of the solidified fiber is at least 0.

40.

18. The method according to claim 17, wherein, The Herman orientation factor of the solidified fiber is at least 0.

42.

19. The method according to any one of claims 1-7, wherein, The Herman orientation factor of the resulting polymer fiber is at least 0.08 greater than that of the first drawn fiber.

20. The method according to claim 19, wherein, The Herman orientation factor of the resulting polymer fiber is at least 0.1 greater than that of the first drawn fiber.

21. The method according to any one of claims 1-7, wherein, The polymer fibers produced have a β c The activation energy of the structural relaxation is less than 700 kJ / mol.

22. The method according to claim 21, wherein, The β of the polymer fiber produced c The activation energy for structural relaxation is less than 650 kJ / mol.

23. The method according to claim 22, wherein, The β of the polymer fiber produced c The activation energy for structural relaxation is less than 600 kJ / mol.

24. The method according to any one of claims 1-7, wherein, The β of the polymer fiber produced c The activation energy for structural relaxation ranges from 500 to 600 kJ / mol.

25. The method according to claim 24, wherein, The β of the polymer fiber produced c The activation energy for structural relaxation ranges from 530 to 570 kJ / mol.

26. The method according to any one of claims 1-7, wherein, The cyclization activation energy of the resulting polymer fiber is at least 7 kJ / mol greater than that of the first drawn fiber.

27. The method according to claim 26, wherein, The cyclization activation energy of the resulting polymer fiber is at least 11 kJ / mol greater than that of the first drawn fiber.

28. The method according to claim 27, wherein, The cyclization activation energy of the resulting polymer fiber is at least 13 kJ / mol greater than that of the first drawn fiber.

29. The method according to any one of claims 1-7, wherein, The resulting polymer fiber has a toughness of at least 4 g / d.

30. The method according to claim 29, wherein, The resulting polymer fiber has a toughness of at least 5 g / d.

31. The method according to claim 30, wherein, The resulting polymer fiber has a toughness of at least 6 g / d.

32. The method according to any one of claims 1-7, wherein, The resulting polymer fiber has a Young's modulus of at least 95 g / d.

33. The method according to claim 32, wherein, The resulting polymer fiber has a Young's modulus of at least 100 g / d.

34. The method according to any one of claims 1-5, wherein, The resulting polymer fiber has a Young's modulus of 95 to 130 g / d.

35. The method according to claim 33, wherein, The Young's modulus of the resulting polymer fiber ranges from 100 to 130 g / d.

36. The method according to claim 34, wherein, The Young's modulus of the resulting polymer fiber ranges from 115 to 125 g / d.

37. The method according to any one of claims 1-7, wherein, The peak temperature of the cyclization exothermic reaction of the resulting polymer fiber is at least 3°C ​​higher than the peak temperature of the cyclization exothermic reaction of the solidified fiber and / or the first drawn fiber.

38. The method according to any one of claims 1-7, wherein, The resulting polymer fiber is a polyacrylonitrile-based polymer fiber.

39. The method according to any one of claims 1-7, wherein, The spinning of this polymer solution is achieved through wet spinning.

40. The method according to any one of claims 1-7, wherein, The coagulation bath contains a mixture of DMSO and water.

41. The method according to any one of claims 1-7, wherein, The resulting polymer fiber is a carbon fiber precursor fiber.

42. A polymer fiber produced by the method according to any one of claims 1-41.

43. A method for producing carbon fiber, the method comprising oxidizing the polymer fiber according to claim 42 or the polymer fiber produced by the method according to any one of claims 1-41 to form a stabilized carbon fiber precursor fiber, and then carbonizing the stabilized carbon fiber precursor fiber to produce carbon fiber.

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