A PROCESS FOR PRODUCING POLYMER FIBER AND POLYMER FIBER MADE THEREFROM

Abstract

The present disclosure relates generally to a process for producing polymer fibers, typically polyacrylonitrile-based fibers, the properties of which are controlled by certain parameters of the process, such as the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch employed. The present disclosure also relates to a process for producing carbon fiber from such polymer fibers.

Description

CROSS REFERENCE TO RELATED APPLICATIONS

This present application claims priority to U.S. provisional application No. 63/157,111, filed Mar. 5, 2021, the entire contents of which is hereby incorporated by reference.

FIELD OF THE INVENTION

The present disclosure relates generally to a process for producing polymer fibers, typically polyacrylonitrile-based fibers, the properties of which are controlled by certain parameters of the process, such as the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch employed. The present disclosure also relates to a process for producing carbon fiber from such polymer fibers.

BACKGROUND

Carbon fibers, particularly polyacrylonitrile (PAN)-based carbon fibers, are increasing in demand across many market sectors due to their high performance-to-weight ratio.¹ However, market ubiquity and wider industrial acceptance are inhibited by their relatively high cost.^2,3 To reduce cost, carbon fiber manufacturers are looking toward larger tow sizes due to their increased capacity, higher throughput, and reduced capital cost per pound of fiber processed.^4-7 While large tows (>24,000 filaments) offer evident cost advantages, they are encumbered by processing challenges resulting in wet spinning being the preferred fabrication method.^8,9 Large tows are plagued by additional downstream challenges including a higher percentage of filament breaks in hot draw,⁸ inability to apply steam stretch or traditional winding mechanisms,^9,10 higher exotherm during thermo-oxidative stabilization (TOS), and variability in carbon fiber yield.¹¹ Despite these additional challenges in large tow processing, property targets and performance metrics have not waned from their small tow counterparts.

To attain carbon fiber property targets, focus has been turned to precursor development and fiber spinning, which establishes the structure and morphology of the carbon fiber at the nascent stages.^12-15 Precursor structure is influenced by chemical composition,^16,17 fiber denier and crystallinity,¹⁸ fibril structure and alignment,^19-21 and spinning speeds.²⁰ Structural imperfections imposed by the spinning process, such as cavities, cracks, and flaws, can impact fiber shrinkage during TOS and reduce the strength of resulting carbon fibers.^22-24 Therefore, elucidating key characteristics of wet spinning on the development of PAN precursor structure can lead to improvement of mechanical properties of carbon fibers.

The morphology and structure of PAN has been well defined in literature,¹³ but researchers were originally riddled by differences in experimental observations under varying process histories.^25-28 Bashir et al. performed a series of elegant studies and clarified that the structure of atactic PAN can be divided into unoriented and oriented states.^29-32 In the unoriented state, PAN adopts a two-phase morphology represented by an amorphous domain and an ordered mesophase domain observed via two thermal transitions in differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA).^39,32 The lower temperature DMA transition, (β_c) c.a. 100° C., is attributed to the mesophase domain, whereas the higher temperature transition, (α) c.a. 130-160° C., is attributed to the amorphous domain. As the PAN fiber becomes oriented on drawing and chain packing occurs, the intensity of the α peak decreases until it vanishes, at which point the oriented PAN becomes a “laterally ordered single-phase”.^26,31,33,34

Beyond thermal transitions and chain packing, chain conformation is also shown to shift as PAN transitions between unoriented and oriented states. Sawai et al. and others have demonstrated that unoriented atactic PAN adopts a greater ratio of irregular helical sequences and, upon drawing, the oriented chains are rearranged into planar zigzag conformations.^35-37 Shen et al. declared syndiotactic systems, with the highest degree of planar zigzag arrangements, to have the lowest energy, highest stiffness, and most favorable chain conformation for inter-molecular packing and TOS.³⁸

While precursor structure has been well defined in the unoriented and oriented states, the evolution of microstructure in PAN-based precursors, or directly connecting the path from unoriented to oriented states in fiber spinning, is rarely studied because mimicking commercial spinning processes and investigating the impact of process changes on fiber structure is capital intensive and difficult. Rather, many studies have focused efforts on nascent stages of fiber spinning where coagulation parameters can have immediate and profound impacts on the precursor structure.^39-44 In particular, much focus has been placed on the dope extrusion rate and jet stretch, or stretch taken between the spinneret and first godet out of coagulation in wet spinning.^45-49

Aside from jet stretch, wet spinning comprises other stages of draw which can be used interchangeably to achieve a target filament denier prior to TOS and carbonization. Stretch may be applied immediately following coagulation in a gel bath or some non-solvent medium,^39,50,51 conventional dry tensile draw,^37,52,53 or through a hot draw step with hot liquid or a steam chest.⁵¹ Although each drawing stage reduces fiber diameter, the mechanism of draw may be vastly different due to the nature of the structure in the gel state vs. the hot drawn state. Edrington's thesis work demonstrated a total draw limit for the combination of draw stages, but did not go as far as decoupling jet stretch and gel draw from hot draw to determine where the greatest impact is on crystal alignment and orientation.⁵¹

Therefore, there is an ongoing need for achieving higher orientated states via optimum spinning processes that could lead to more stable crystalline regions and an advantageous precursor structure for the conversion to carbon fiber. Herein, a process for the production of polymer fiber, such fiber being suitable for formation of carbon fiber, is described. The inventive process employs understandings derived from investigation of microstructure with wide angle x-ray scattering (WAXS) along with thermal transitions in DMA and stabilization kinetics via DSC of PAN-based polymer fibers.

SUMMARY OF THE INVENTION

Advantageously, it has surprisingly been discovered that given stages of stretch are not reciprocal and that stretch profile may play a significant role in precursor structure. The inventive process employs understandings derived from investigation of microstructure with wide angle x-ray scattering (WAXS) along with thermal transitions in DMA and stabilization kinetics via DSC of polymer fibers, such as PAN-based fibers.

In a first aspect, the present disclosure relates to a process for producing polymer fiber, the process comprising:

• • a) spinning a polymer solution into a coagulation bath, wherein a jet stretch is applied, to form coagulated fiber;
  • b) subjecting the coagulated fiber obtained in step (a) to a wet stretch to form a first drawn fiber; and
  • c) subjecting the first drawn fiber obtained in step (b) to a hot stretch, thereby forming the polymer fiber;
  • wherein the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch are effective to achieve the following properties wherein:
    • the Herman's orientation factor of the polymer fiber produced is at least 0.60, typically at least 0.65, more typically at least 0.67, and
    • the crystallite thickness of the polymer fiber produced 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.

In a second aspect, the present disclosure relates to a process for producing carbon fiber, the process comprising oxidizing the polymer fiber described herein or the polymer fiber produced according to the process described herein to form stabilized carbon fiber precursor fibers and then carbonizing the stabilized carbon fiber precursor fiber, thereby producing carbon fiber.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 shows the structural evolution, as observed by WAXS analysis, of the fibers throughout the process described herein in which a certain stretch profile was used.

DETAILED DESCRIPTION

As used herein, the terms “a”, “an”, or “the” means “one or more” or “at least one” and may be used interchangeably, unless otherwise stated.

As used herein, the term “and/or” used in a phrase in the form of “A and/or B” means A alone, B alone, or A and B together.

As used herein, the term “comprises” includes “consists essentially of” and “consists of.” The term “comprising” includes “consisting essentially of” and “consisting of.” “Comprising”, which is synonymous with “including,” “containing,” or “characterized by,” is intended to be inclusive or open-ended and does not exclude additional, unrecited elements or steps. The transitional phrase “consisting essentially of” is inclusive of the specified materials or steps and those that do not materially affect the basic characteristic or function of the composition, process, method, or article of manufacture described. The transitional phrase “consisting of” excludes any element, step, or component not specified.

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

As used herein, and unless otherwise indicated, the term “about” or “approximately” means 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 certain embodiments, the term “about” or “approximately” means within 1, 2, 3, or 4 standard deviations. In certain embodiments, the term “about” or “approximately” means 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.

Also, it should be understood that any numerical range recited herein is intended to include all sub-ranges subsumed therein. For example, a range of “1 to 10” is intended to include all sub-ranges between and including the recited minimum value of 1 and the recited maximum value of 10; that is, having a minimum value equal to or greater than 1 and a maximum value of equal to or less than 10. Because the disclosed numerical ranges are continuous, they include every value between the minimum and maximum values. Unless expressly indicated otherwise, the various numerical ranges specified in this application are approximations.

Throughout the present disclosure, various publications may be incorporated by reference. Should the meaning of any language in such publications incorporated by reference conflict with the meaning of the language of the present disclosure, the meaning of the language of the present disclosure shall take precedence, unless otherwise indicated.

The first aspect of the present disclosure relates to a process for producing polymer fiber, the process comprising:

• • a) spinning a polymer solution into a coagulation bath, wherein a jet stretch is applied, to form coagulated fiber;
  • b) subjecting the coagulated fiber obtained in step (a) to a wet stretch to form a first drawn fiber; and
  • c) subjecting the first drawn fiber obtained in step (b) to a hot stretch, thereby forming the polymer fiber;
  • wherein the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch are effective to achieve the following properties wherein:
    • the Herman's orientation factor of the polymer fiber produced is at least 0.60, typically at least 0.65, more typically at least 0.67, and
    • the crystallite thickness of the polymer fiber produced 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.

In step a) of the process, a polymer solution, or “spin dope”, is spun into a coagulation bath to form a coagulated fiber. Typically, the polymer solution is a homogeneous solution comprising a polyacrylonitrile-based polymer and a solvent. Accordingly, in an embodiment, the polymer fiber produced is polyacrylonitrile-based polymer fiber.

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

The polyacrylonitrile-based polymer comprises 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.

In an embodiment, 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-ethylhexylacrylate, isopropyl acetate, vinyl acetate (VA), vinyl propionate, vinyl imidazole (VIM), acrylamide (AAm), diacetone acrylamide (DAAm), allyl chloride, vinyl bromide, vinyl chloride, vinylidene chloride, sodium vinyl sulfonate, sodium p-styrene sulfonate (SSS), sodium methallyl sulfonate (SMS), sodium-2-acrylamido-2-methyl propane sulfonate (SAMPS), and mixtures thereof.

The comonomer ratio (amount of one or more comonomers to amount of acrylonitrile) is not particularly limited. However, a suitable comonomer ratio is 0 to 20%, typically 1 to 5%, more typically 1 to 3%.

The molecular weight of the polyacrylonitrile-based polymers suitable for use according to the described process may be within the range of 60 to 500 kg/mole, typically 90 to 250 kg/mole, more typically 115 to 180 kg/mole.

Suitable solvents for the polymer may be selected from the group consisting of dimethyl sulfoxide (DMSO), dimethyl formamide (DMF), dimethyl acetamide (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), dimethyl formamide (DMF), dimethyl acetamide (DMAc), ethylene carbonate (EC), N-methyl-2-pyrrolidone (NMP).

The polymer solution used is typically free of gels and/or agglomerated polymer. The presence of gels and/or agglomerated polymer may be determined using any method known to those of ordinary skill in the art. For example, a Hegman gauge may be used to determine the presence of gels and/or agglomerated polymer. The polymer solutions used are generally stable and do not exhibit gel formation over time.

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

The polymer solution may be subjected to conventional wet spinning and/or air-gap spinning after removing air bubbles by vacuum. In wet spinning, the dope is filtered and extruded through holes of a spinneret (typically made of metal) into a liquid coagulation bath for the polymer to form filaments. The spinneret holes determine the desired filament count of the fiber (e.g., 3,000 holes for 3K carbon fiber). 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 coagulating bath. In an embodiment, the spinning of the polymer solution is achieved by wet spinning.

The coagulation liquid used in the process is a mixture of solvent and non-solvent. Water or alcohol is typically used as the non-solvent. Suitable solvents include the solvents described herein. In an embodiment, dimethyl sulfoxide, dimethyl formamide, dimethyl acetamide, or mixtures thereof, is used as solvent. In another embodiment, dimethyl sulfoxide is used as solvent. The ratio of solvent and non-solvent, and bath temperature are not particularly limited and may be adjusted according to known methods to achieve the desired solidification rate of the extruded nascent filaments in coagulation. However, the coagulation bath typically comprises 40 wt % to 85 wt % of one or more solvents, the balance being non-solvent. In an embodiment, the coagulation bath comprises a mixture of DMSO and water.

A jet stretch is applied in step a). As used herein, the amount of jet stretch applied in the coagulation step is the ratio of the first roller take-up velocity to the dope extrusion velocity as would be understand by a person of ordinary-skill in the art.^46,48 Dope extrusion velocity, V_jet, is calculated by Equation 1 below, where Q is the volumetric flow rate, typically provided by a metering pump, N is the number of holes in the spinneret, and D is the diameter of each hole. Accordingly, a person of ordinary skill in the art would recognize that V_jet may be adjusted by selecting suitable values for Q, N, and D.

$\begin{array}{cc}{V}_{jet}=\frac{4Q}{\pi N{D}^2}& \left(1\right)\end{array}$

In step b) of the process, coagulated fiber obtained in step (a) is subjected to a wet stretch to form a first drawn fiber. As used herein, wet stretch is synonymous with “wet draw”, “first draw”, “1^st draw”, and “FD” and the terms may be used interchangeably. During wet stretch, the coagulated fiber is immersed in and conveyed through one or more baths, typically through the action of rollers or stretch rolls. The amount of wet stretch, as used herein, is defined as the ratio of the speed of the stretch rolls between coagulation and the 1st draw. The one or more baths used for wet stretch may each be maintained at a temperature 40° C. to 100° C.

In step c) of the process, the first drawn fiber obtained in step (b) is subjected to a hot stretch, thereby forming the polymer fiber. As used herein, hot stretch is synonymous with “hot draw”, “second draw”, “2^nd draw”, and “HD” and the terms may be used interchangeably. During hot stretch, the first drawn fiber is conveyed through a heat source without immersion in liquid, typically through the action of rollers or stretch rolls. For example, the heat source may be steam or a tubular furnace through which the first drawn fiber obtained in step (b) is conveyed. The amount of hot stretch, as used herein, is defined as the ratio between the stretching rolls following FD and the final winder speed.

The process of the present disclosure may be conducted continuously or in a batch manner. As used herein, a process “conducted continuously” refers to a process in which the fiber is conveyed through one or more processing steps a single work unit at a time without any breaks in time, substance, or sequence. This is in contrast to a batch process, which would be understood as being a process that comprises a sequence of one or more steps that are performed in a defined order and in which a finite quantity of material is treated or produced at the end of the sequence, which must be repeated in order to treat or produce another batch of material. In an embodiment, the process is conducted continuously.

The polymer fiber, typically polyacrylonitrile-based fiber, produced by the process described herein may be employed as a precursor fiber, so-called white fiber (“WF”), for the production of carbon fiber. Thus, in an embodiment, the polymer fiber produced is carbon fiber precursor fiber.

In the process described herein, the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch are effective to achieve the following properties, namely the Herman's orientation factor of the polymer fiber produced is at least 0.60, typically at least 0.65, more typically at least 0.67, and the crystallite thickness of the polymer fiber produced 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.

It has been found that coagulation influences initial structure and dictates orientation and structural evolution, that hot draw, or high temperature draw, post-wet stretch is beneficial for increasing the size of the mesophase domain, and that structural order directly impacts mechanical properties and downstream process parameters for TOS. Thus, the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch are adjusted such that certain properties are achieved.

Wide Angle X-Ray Scattering (WAXS) is used to measure parameters that are then used to determine various properties of the fibers throughout the process, including Herman's orientation factor, crystallite thickness, crystallinity, and D-spacing.

WAXS measurements are performed according to methods known to those of ordinary skill in the art. In a suitable example, WAXS is performed in transmission mode using a Xenocs Xeuss 2.0 system. The source is GeniX3DCu ULD 8 keV with wavelength, λ, of 1.54189 Å. A single fiber tow (1,000 filaments) is aligned and fixed across an aperture card. The aperture card with the aligned fiber bundle is then transferred to the WAXS sample holder, which is placed at 101.17 mm from the 2D detector (Dectris Pilatus 200 K). Exposure time for precursor fibers is 600 s. Data processing to obtain the integrated diffracted intensity versus 2θ and azimuthal angle, φ, may be performed using the software Foxtrot provided by Xenocs. Structural parameters may be determined by peak fitting of the Foxtrot data using Origin 2017 software using a Lorentzian function with two crystalline peaks located at approximately 2θ₁≈16.9° and 2θ₂≈29.3° and an amorphous peak centered about 2θ_a≈25°.⁵⁴

Crystallinity, X_c, is determined by Equation 2, where A_θ1 and A_θ2 are the areas of the two crystalline peaks, 2θ₁ and 2θ₂ respectively, and A_θa is the area of the amorphous peak. D-spacing is determined by Bragg's Law (Equation 3) and crystallite thickness, L_c, by the Scherrer Equation (Equation 4), where the shape factor, K, is 0.89, and β is the full width half maximum in radians of the corresponding crystalline peaks. Azimuthal scan profiles are acquired by integrating the diffraction intensities, I, azimuthally at the (100) diffraction plane (2θ₁≈16.9°).^55,56 Herman's orientation factor (f_c) is determined by Equation 5 and 6.³⁹

$\begin{array}{cc}{\chi }_c=\frac{A_{{\theta }_1}+A_{{\theta }_2}}{A_{{\theta }_1}+A_{{\theta }_2}+A_{{\theta }_a}}& \left(2\right)\end{array}$ $\begin{array}{cc}d=\frac{\lambda }{2\sin \theta }& \left(3\right)\end{array}$ $\begin{array}{cc}{L}_c=\frac{K\lambda }{\mathrm{\beta cos\theta }}& \left(4\right)\end{array}$ $\begin{array}{cc}\langle \cos \phi \rangle =\frac{{\int }_0^{\pi }I\left(\phi \right){\cos }^2\mathrm{\phi sin\phi }d\phi }{{\int }_0^{\pi }I\left(\phi \right)\sin \phi d\phi }& \left(5\right)\end{array}$ $\begin{array}{cc}{f}_c=\frac{3\langle {\cos }^2\vartheta \rangle -1}{2}& \left(6\right)\end{array}$

The Herman's orientation factor, crystallite thickness, crystallinity, and D-spacing may be determined at any point in the process described herein. For example, fiber samples may be taken after each step of the process and measured by WAXS to determine the Herman's orientation factor, crystallite thickness, crystallinity, and D-spacing of the fibers formed in each step.

In an embodiment, the Herman's orientation factor of the polymer fiber produced is at least 0.60, typically at least 0.65, more typically at least 0.67.

In another embodiment, the Herman's orientation factor of the coagulated fiber is at least 0.35, typically at least 0.40, more typically at least 0.42.

In yet another embodiment, the Herman's orientation factor of the polymer fiber produced is at least 0.08, typically at least 0.1, greater than the Herman's orientation factor of the first drawn fiber.

In an embodiment, the crystallite thickness of the polymer fiber produced 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.

In an embodiment, the crystallinity of the coagulated fiber is no more than 8%, typically no more than 7%, more typically no more than 6%, greater than the crystallinity of the first drawn fiber.

Throughout the process, the fibers obtained in each step may be characterized by their cross-sectional diameters. Any suitable method of measuring the diameters of the fibers may be used. In a suitable example, Scanning Electron Microscopy (SEM) is used. In preparation for imaging, the polymer fibers are freeze-dried on a Labconco freeze dryer using a tert-butanol/water cosolvent system. Filament cross sections are produced by saturating the fiber tow in water and then submerging it in liquid nitrogen before fracturing. Filament ends are mounted on 15 mm aluminum SEM stubs using carbon tape and sputter-coated with 3 nm of platinum to reduce charging. A Hitachi S-4800 may be used to image the cross sections at 2 kV and varied magnifications. Average filament diameters may be measured using QUARTZ PCI image processing software. A minimum of 10 measurements are taken per sample.

In an embodiment, the average diameter of the coagulated fiber is at least 40 μm, typically at least 45 μm, more typically at least 50 μm, still more typically at least 55 μm.

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

The fibers obtained in each step of the process may also be characterized by an activation energy, ΔH, for the structural relaxation at the mesophase glass transition temperature, β_c.

The activation energy, ΔH, for the structural relaxation at the mesophase glass transition temperature, β_c, is measured using Dynamic Mechanical Analysis (DMA) using methods known to those of ordinary skill in the art. Suitably, DMA is performed on a TA Instruments Discovery Hybrid Series HR-2 Rheometer. Precursor fiber tows are tested using a rectangular tension fixture geometry in tension mode. A 5-minute temperature soak and 1±0.1 N of axial force is applied to condition the samples. The loss modulus, storage modulus, and tan δ variables are plotted against temperature during an oscillation temperature ramp at varied frequencies of 0.5, 1, 5, 10, and 15 Hz, from a temperature range of 45° C. to 165° C. at 1° C./min. Axial strain is set to 0.1%. The activation energy, ΔH, for the structural relaxation at the mesophase glass transition temperature, β_c, taken as the peak in tan δ, is calculated by an Arrhenius relationship with frequency, f, and the universal gas constant, R, by Equation 7.^33,35,57

$\begin{array}{cc}\Delta H=-R\frac{a\left(\ln f\right)}{d\left(\frac{1}{{\beta }_c}\right)}& \left(7\right)\end{array}$

In an embodiment, the activation energy of structural relaxation for β_c of the polymer fiber produced is less than 700 kJ/mol, typically less than 650 kJ/mol, more typically less than 600 kJ/mol.

In another embodiment, the activation energy of structural relaxation for βc of the polymer fiber produced is from 500 to 600 kJ/mol, typically from 530 to 570 kJ/mol.

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

Cyclization activation energy is determined from the Flynn-Wall-Ozawa (FWO) method, Equation 8, where E_a is the cyclization activation energy, R is the universal gas constant, T_pk is the peak exotherm temperature in Kelvin, and ω is the temperature ramp rate in Kelvin.^58,59

$\begin{array}{cc}-\frac{E_a}{R}=0.951\times \frac{d\left(\ln \omega \right)}{d\left(\frac{1}{T_{pk}}\right)}& \left(8\right)\end{array}$

In an embodiment, the cyclization activation energy of the polymer fiber produced is at least 7 kJ/mol, typically at least 11 kJ/mol, more typically at least 13 kJ/mol, greater than the cyclization activation energy of the first drawn fiber.

Also using DSC, the peak temperature of the cyclization exotherm of the fibers obtained in each step of the process may be determined. The cyclization reaction is generally exothermic and may be viewed as a peak in a DSC scan in which heat flow, in units of W/g, is plotted as a function of temperature. In an embodiment, the peak temperature of the cyclization exotherm of the polymer fiber produced is at least 3° C. higher than the peak temperature of the cyclization exotherm of the coagulated fiber and/or the first drawn fiber.

The polymer fiber formed by the inventive process described herein, typically carbon fiber precursor fiber, or white fiber, have mechanical properties, such as tenacity, elongation, and Young's modulus, that result from the surprising observation that stretch profile, i.e., the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch, give rise to certain microstructure.

Linear mass density, or tow denier, of the polymer fiber formed may be calculated from the average of two yield measurements based on the weight of a 3-meter sample section. The linear mass density may be expressed in terms of a tow (bundle of filaments) or on a per filament basis. Herein, the linear mass density is expressed on a per filament basis. Thus, linear mass density used herein is the mass in grams per 9000 meters of a filament, or denier per filament. The linear mass density of the polymer fiber produced is from 0.7 to 1.2, typically from 0.85 to 1.0 denier per filament.

The mechanical properties may be tested using known methods. For example, MTS Criterion C43 and Testworks 4 software may be used. Fiber samples are loaded at a gauge length of 7.875 inches with a preload force of 0.1 lbf and a grip pressure of 60 psi. Each sample is tested a total of 3 times at a cross head speed of 2 in/min.

In an embodiment, the tenacity of the polymer fiber produced is at least 4 g/d, typically at least 5 g/d, more typically at least 6 g/d.

In an embodiment, the Young's Modulus of the polymer fiber produced is at least 95 g/d, more typically at least 100 g/d.

In another embodiment, the Young's Modulus of the polymer fiber produced is 95 to 130 g/d, typically from 100 to 130 g/d, more typically from 115 to 125 g/d.

In the second aspect, the present disclosure relates to a process for producing carbon fiber, the process comprising oxidizing the polymer fiber described herein or the polymer fiber produced according to the process described herein to form stabilized carbon fiber precursor fibers and then carbonizing the stabilized carbon fiber precursor fiber, thereby producing carbon fiber.

The polymer fiber formed according to the process described herein may be oxidized to form stabilized carbon fiber precursor fibers and, subsequently, the stabilized carbon fiber precursor fiber are carbonized to produce carbon fibers.

During the oxidation stage, the polymer fiber are fed under tension through one or more specialized ovens, each having a temperature from 150 to 300° C., typically from 200 to 280° C., in which heated air is fed into each of the ovens. The oxidation process combines oxygen molecules from the air with the fiber and causes the polymer chains to start crosslinking, thereby increasing the fiber density to suitable levels. Such oxidized, typically PAN, fiber has an infusible ladder aromatic molecular structure and it is ready for carbonization treatment.

Carbonization results in the crystallization of carbon molecules and consequently produces a finished carbon fiber that has more than 90 percent carbon content. Carbonization of the oxidized, or stabilized, carbon fiber precursor fibers occurs in an inert (oxygen-free) atmosphere, typically nitrogen atmosphere, inside one or more specially designed furnaces. The oxidized carbon fiber precursor fibers are passed through one or more ovens each heated to a temperature of from 300° C. to 1650° C.

Surface treatment may include pulling the carbonized fiber through an electrolytic bath containing an electrolyte, such as ammonium bicarbonate or sodium hypochlorite. The chemicals of the electrolytic bath etch or roughen the surface of the fiber, thereby increasing the surface area available for interfacial fiber/matrix bonding and adding reactive chemical groups useful for forming composites.

Next, the carbon fiber may be subjected to sizing, where a size coating, e.g. epoxy-based coating, is applied onto the fiber. Sizing may be carried out by passing the fiber through a size bath containing a liquid coating material. Sizing protects the carbon fiber during handling and processing into intermediate forms, such as dry fabric and prepreg. Sizing also holds filaments together in individual tows to reduce fuzz, improve processability and increase interfacial shear strength between the fiber and the matrix resin used in manufacturing composite materials. Following sizing, the coated carbon fiber is dried and then wound onto a bobbin. The carbon fiber produced are suitable for use in the production of composite materials.

The process according to the present disclosure and polymer fibers produced therefrom are further illustrated by the following non-limiting examples.

EXAMPLES

Example 1. Preparation of Polymer Fiber

Polymer dope was prepared by mixing the PAN-based polymer in DMSO in a 15-gallon Myers Mixer at a target concentration of 18.5 wt. % solids. The polymer dope was warmed to 65° C., deaerated 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 prior to wet spinning. The dope flow rate was controlled at a constant rate of 40 mL/min and extruded through a 1,000-hole spinneret (hole diameter=50 μm) into a heated coagulation bath containing a mixture of DMSO and water.

Polymer composition, bath concentrations, temperatures, and other spinning parameters remained constant while the only variable was the ratio of stretch taken during jet stretch and 2^nd draw. Three variants were studied with a proportionate amount of stretch taken between jet stretch and 2^nd draw as outlined in Table 1 below in which the variant is designated as “1”, “2”, or “3” and the step from which the fiber sample was taken is designated as “Coag” (following coagulation), “FD” (following 1^st draw), and “WF” (following 2^nd draw). The amount of jet stretch applied in coagulation was controlled by the ratio, X, of the first roller take-up velocity to the dope extrusion velocity.^46,48 Dope extrusion velocity, V_jet, is calculated by Equation 1 described hereinabove. The ratio of speeds in the stretch rolls between coagulation (“Coag”) and 1^st draw (“FD”) is the amount of wet stretch, Y, and the ratio between the stretching rolls following FD and the final winder speed is the amount of stretch applied in hot draw, Z.

For Coag-2 and Coag-3 the jet stretch was increased by 36% and 58% in relation to Coag-1, respectively, and 2^nd draw was reduced by a commensurate amount. In all three variants studied, the final winder speed was equivalent and, therefore, the total stretch applied for each white fiber precursor was equivalent, X*Y*Z. For each of the three variant samples collected following coagulation and 1^st draw were allowed to hang dry to constant weight before further analysis. White fiber (WF) samples were collected at the winder following a drying process during spinning and used as is.

TABLE 1
Samples collected along spin line with various stretch ratios.
				1st Drawb)	2nd Drawc)
Sample		Sampling	Jet Stretcha)	(Wet	(Hot	Total
ID	Profile	Location	(Coagulation)	Stretch)	Stretch)	Stretchd)
Coag-1	1	Coagulation	X	—	—	X
FD-1		1st Draw	X	Y	—	X*Y
WF-1		Winder	X	Y	Z	X*Y*Z
Coag-2	2	Coagulation	1.36*X	—	—	1.36*X
FD-2		1st Draw	1.36*X	Y	—	1.36*X*Y
WF-2		Winder	1.36*X	Y	0.74*Z	X*Y*Z
Coag-3	3	Coagulation	1.58*X	—	—	1.58*X
FD-3		1st Draw	1.58*X	Y	—	1.58*X*Y
WF-3		Winder	1.58*X	Y	0.63*Z	X*Y*Z
a)ratio of 1st stretching rolls/dope extrusion velocity;
b)ratio of 2nd stretching rolls/1st stretching rolls;
c)ratio of winder speed/2nd stretching rolls;
d)total stretch: jet stretch*1st draw*2nd draw

Example 2. SEM Analysis

SEM analysis, as described hereinabove, was conducted on fiber samples indicated in Table 1.

It was observed that as jet stretch increases from Coag-1 to Coag-2 to Coag-3, filament cross-sections transition from a kidney-bean shape to an oblong round shape and, finally, a more circular shape. Jet stretch is known to influence the shear forces in the spinneret capillary, which decreases die swell and, thereby, decreases the solvent exchange.^46,48,60 Increasing jet stretch also decreases filament diameter as it exits the capillary, wherein a greater protofiber radius has a thinner and sparser skin structure making counter diffusion of water and DMSO easier.⁶⁰ The larger diameter gives Coag-1 a milder skin-core transition as compared to Coag-2 and Coag-3, which have more distinct skin layers. For Coag-3, counter diffusion is hindered enough to form a solid skin and impede complete coagulation of the core structure leaving a looser sponge-like character as compared to Coag-1 and Coag-2. In addition, the higher jet stretch is confounded by a shorter residence time in the coagulation bath, which may also leave the fiber not completely coagulated.

Following coagulation, the fibers were all subjected to the same draw ratios in first draw, albeit the actual speeds and residence times varied due to differences in jet stretch. The cross-sectional shape converges toward a round filament in all three cases, but FD-1 retains some kidney-bean resemblance. Cross-sectional images were obtained by fracturing the samples in liquid nitrogen where fracture surfaces can give insight into defects.⁶¹ FD-1 and FD-2 show radiating patterns from the center of the fiber which is an indication of an internal flaw. FD-3 shows a separation of the skin layer from the core suggesting the fracture initiated in the skin and propagated around the core of the filament rather than through the center of the filament. The filament skins also show apparent differences with FD-1 having finer striations along the fiber axis whereas FD-3 has more pronounced and deeper skin grooves due to the faster coagulation rate.⁶² The average filament diameters retain the same proportional difference following first draw as with coagulation, which is expected for a constant first draw ratio. For instance, Coag-2 and FD-2 average filament diameters (49.6 μm and 19.1 μm) remain about 86% of Coag-1 and FD-1 (58.0 μm and 22.2 μm), respectively. Proportionate adjustments in 2^nd draw are used to compensate for speed differences such that the winder speed for each of the three profiles was equal. The average diameters of the coagulated fibers and first drawn fibers as observed through SEM are summarized in Table 2 below. Morphological testing of WF samples via SEM is difficult due to the ductile nature of the fiber and as such is not included.

TABLE 2
Sample ID Average diameter (μm)
Coag-1    58.0
Coag-2    49.6
Coag-3    44.0
FD-1      22.2
FD-2      19.1
FD-3      17.5

Example 3. WAXS Analysis

To investigate the structural differences of the fibers following each stage of stretch, WAXS was performed on each sample. FIG. 1 shows the structural evolution of Profile 1, wherein the scans for each sampling location are overlaid. Interestingly, Coag-1 displays a strong peak c.a. 2θ₁≈17° corresponding to the mesophase ordered region in PAN,³¹ which indicates a significant amount of crystallization occurs as a result of coagulation and jet stretch. Coag-1 also reveals a broad amorphous peak c.a. 2θ_a≈25° and the onset of another higher order diffraction plane c.a. 2θ₂≈29°.⁵⁴ As the fiber is further stretched in first draw, FD-1 shows slight narrowing of 2θ₁ and a greater intensity of the 2θ₂ peak as compared to Coag-1. Following 2^nd draw, WF-1 reveals a much sharper 2θ₁ diffraction peak, very pronounced 2θ₂ peak, and major reduction of the relative 2θ_a intensity.

Crystallinity (X_c), crystallite thickness (L_c), and orientation (f_c) under all processing conditions were determined as described hereinabove. Crystallinity appears to decrease through the process; however this behavior is believed to be an anomaly due to the data processing technique used. Interpretation of crystallinity is often disputed in the literature,¹³ while the suggested method used herein accounts for both 2θ₁ and 2θ₂ peaks as part of the ordered domain as observed in WAXS peak parameters and peak fits for the various fiber samples (not shown). Nonetheless, difficulty in deconvolution of 2θ₂ with the amorphous region, 2θ_a, leads to higher apparent crystallinities upon coagulation, meaning samples with higher amorphous content can artificially increase the peak area for the 2θ₂ peak and thus the total crystallinity of the sample.

Surprisingly, a prominent 2θ₁ peak for coagulation samples indicates that the mesophase domain is largely established at the onset of spinning. While crystallinity and d-spacing for 2θ₁ (d_2θ1) remain relatively stable throughout the stretching process in spinning, significant structural changes occur in crystallite thickness and orientation. Orientation gradually increases from Coag to FD to WF for all three samples with narrowing peaks for each respective azimuthal scan. Crystallite thickness has a small increase between Coag to FD and a surge from FD to WF noted by the narrowing of the 2θ₁ peak. This suggests that hot draw is more effective than wet stretch at fusing the ordered regions together and merging amorphous and ordered domains. Liu et al. have suggested that initial stages of stretch are spent through straightening and disentangling of polymer chains,²² which may also explain the marginal increase in L_c during wet stretch. The WAXS parameters for the fiber samples are summarized in Table 3 below.

TABLE 3
WAXS parameters for fiber samples.
	Sample		χc	2θ1	d2θ1	Lc
	ID	fc a)	[%]b)	[degrees]	[Å]c)	[nm]d)
	Coag-1	0.44	49.71	16.77	5.28	3.93
	FD-1	0.60	44.58	16.83	5.26	3.99
	WF-1	0.70	36.14	16.77	5.28	8.47
	Coag-2	0.42	41.51	16.79	5.28	3.85
	FD-2	0.57	35.50	16.73	5.30	4.05
	WF-2	0.66	35.49	16.86	5.25	7.92
	Coag-3	0.38	45.29	16.76	5.29	3.84
	FD-3	0.55	37.56	16.69	5.31	3.98
	WF-3	0.63	38.78	16.80	5.27	7.86
	a) fc: Herman's orientation factor by azimuthal analysis about 2θ1 peak;
	b)χc: crystallinity;
	c)d-spacing for 2θ1;
	d)Lc: crystallite thickness for 2θ1

Example 4. Molecular Dynamics Simulation

To visualize the evolving structure suggested from the WAXS data, a molecular dynamics simulation was performed. Here, a PAN system consisting of twenty separate chains was subjected to an axial strain to simulate the stretching process.

Schrodinger Materials Science 19-3 and OPLS3-e force-field was used for all simulations. PAN chains were constructed using the “Crosslink Polymer” module by first placing 5000 acrylonitrile monomers and 20 initiator molecules into a periodic box followed by subsequent 50 ns NPT and NVT stages at 400K to densify and relax the cell. A dummy atom was used to represent the propagating radical. Twenty reactions were allowed per step, to correspond to the number of initiators, however a monomer had to be within 4.0 A distance criteria to the dummy atom for a reaction to occur. After each step the system was relaxed for 50 ps. Polymerization was carried out at 400K and continued until conversion of monomer equaled 100%. Following polymerization the system was annealed for an additional 50 ns NPT and NVT at 400K. Stretching simulation was performed by extending one dimension of the periodic box by 0.1% strain for 600 steps to reach a total of 60% strain. A 0.5 Poisson ratio was applied to control the lateral dimensions of the box. Stretch was performed under NVT conditions. Backbone torsions were analyzed following the stretch simulation where torsion was defined by a 4 atom dihedral angle from four connected backbone atoms.

Snapshots of the system through the simulation where a single chain has been represented with enlarged atoms to track its progression were observed. The 0% system was omitted due to being completely amorphous and not considered a good representation of the actual polymer. Enlarged regions for the 10% and 60% strain systems show the local ordering of the atatic polymer chains where at lower strains, a higher population of gauche configurations act as kinks in the chain, disrupting the local order. Observation of the simulation results highlighted the different populations of backbone dihedral angles where there is a high preference for two configurations, the trans (≈180°) and the gauche (≈60°). As stretch progresses, the gauche conformations transform into trans conformations resulting in an extended planar zig zag arrangement and allowing neighboring chains to align and condense, thus forming larger mesophase domains. As the stretch progresses further, the gauche conformations tend toward zero and the systems approaches a single phase behavior.

Example 5. DMA Analysis of Fibers

The evolution of the storage (E′) and loss (E″) moduli and the corresponding tan delta peaks for the 10 Hz frequency scan for Profile 1, respectively, was investigated using DMA analysis as described hereinabove. Due to the challenges in taking accurate measurements of the cross sectional area of the different samples, the magnitudes of the storage and loss moduli should not be viewed as significant. Rather, tan delta which is a ratio of E″ and E′, will be used to compare the samples as it eliminates the bias of sample dimensions.

Coag-1 displayed two peaks in the tan delta curve, an α transition peak c.a. 145° C., and a β_c transition peak c.a. 105° C.^31,33 As the fiber becomes oriented in FD-1 the α transition disappears and the magnitude of the β_c transition initially increases suggesting the α peak corresponds to the amorphous region and the β_c corresponds to the ordered mesophase region.³¹ Coag-1 also shows a narrower β_c transition peak compared to FD-1 due to the more distinguished and sizable amorphous regions for as-spun fibers.⁶³ As stretch progresses through WF-1, transition toward single phase behavior is further supported with only a single β_c peak being present.

Activation energies for the β_c transition were determined by an Arrhenius fit plotting the reciprocal of peak temperatures,^1/β_c, as a function of logarithm of frequency, In f. The results are summarized in Table 4 below.

TABLE 4
Peak tan δ temperatures at varying frequency (for 1° C./min ramp)
and activation energies ΔH.
Sample	βc peak tan δ temperature [° C.]a)	ΔH
ID	0.5 Hz	1 Hz	5 Hz	10 Hz	15 Hz	[kJ/mol]b)
Coag-1	101.71	103.19	105.00	105.87	107.16	799
FD-1	98.86	100.03	102.92	104.94	105.77	572
WF-1	96.73	99.16	101.14	103.11	104.26	557
Coag-2	102.31	103.21	105.33	106.50	106.77	879
FD-2	100.21	102.87	103.95	105.71	106.13	700
WF-2	99.05	100.50	103.31	104.15	105.71	631
Coag-3	103.24	103.77	106.26	106.74	107.49	927
FD-3	102.71	103.62	106.05	106.70	107.45	867
WF-3	98.47	100.27	102.31	103.70	104.83	664
a)βc: peak relaxation temperature as determined by signal maximum at each given
frequency for temperature ramp of 1° C./min and 0.1% strain;
b)ΔH: activation energy of structural relaxation for βc as determined by Arrhenius relationship,
$\Delta H=-R\frac{d\left(\ln f\right)}{d\left(\frac{1}{{\beta }_c}\right)},$ where f is frequency and R is the universal gas constant

With increasing frequency, the transition temperature increases and magnitude of tan δ decreases due to the shorter response times and viscoelastic nature of the PAN-based fiber. The magnitude of the slope for stretch Profile 1 follows in the order of Coag-1>FD-1>WF-1, which corresponds to decreasing values of ΔH as the fiber is stretched.

Example 6. DSC Analysis

DSC analysis was conducted on the fiber samples obtained from the process described herein. DSC results for the evolution of the fibers at 20° C./min in nitrogen were determined. A distinct onset peak pertaining to the exothermic cyclization reaction was observed for Coag-1 c.a. 243° C., which diminishes in FD-1 and is absent in WF-1. This initial peak is attributed to the larger sized amorphous domains for Coag-1 and FD-1 with respect to WF-1. WF-1 shows a more gradual onset for exotherm, perhaps due to the greater commingling of ordered and amorphous domains as compared to Coag-1 and FD-1, which may have more discrete phase boundaries. With the greater onset peak and higher internal heat generation below 250° C., the cyclization exotherm in the ordered domain also peaks at a lower temperature and higher heat flow for Coag-1 and FD-1 c.a. 296° C. as compared to WF-1 c.a. 299° C.

The activation energy for cyclization, E_a, was determined from the Ozawa method where inverse of peak exotherm temperature, 1/T_pk, is plotted against the logarithm of the heating rate, In ω, which has been published with peak temperatures increasing for faster ramp rates.^59,64 Peak temperatures and calculated E_a values are summarized Table 5 below.

TABLE 5
DSC peak maxima determined in nitrogen at various
ramp rates and cyclization activation energy
	Peak temperature [° C.]a)	Ea
Sample	2° C./	5° C./	10° C./	20° C./	30° C./	[kJ/
ID	min	min	min	min	min	mol]b)
Coag-1	261.41	273.87	284.09	296.04	303.50	156.7
FD-1	261.85	274.30	284.68	296.71	304.14	156.1
WF-1	267.04	278.56	288.33	299.29	306.52	169.8
Coag-2	261.44	273.94	284.45	296.41	303.92	155.2
FD-2	262.19	274.75	285.08	297.10	304.33	156.7
WF-2	265.14	276.73	286.25	297.49	304.78	168.1
Coag-3	261.38	274.01	284.27	295.99	304.02	155.5
FD-3	262.33	274.76	284.93	297.04	304.34	157.2
WF-3	264.12	275.94	285.53	297.12	304.15	164.9
a)Peak exotherm temperature for maximum heat flow in DSC in nitrogen at given ramp rate;
b)Ea: cyclization activation energy determined by Ozawa method

Cyclization activation energy remains nearly constant between samples collected after coagulation and those following first draw with values ranging between 155 kJ/mol to 158 kJ/mol. The most significant change occurs upon hot draw when the E_a value increases to 169.8 kJ/mol in the case of WF-1. Cyclization onset has been attributed to structural abnormalities and defects,^64-66 which may explain why WF samples have much higher activation energies. White fibers following hot draw contain the greatest sized ordered domains as evidenced by WAXS and DMA analysis, which leave fewer structural defects for onset of cyclization.

Example 7. Mechanical Properties of White Fiber

Mechanical properties were used for comparison of the final fibers and are summarized in Table 6 below.

TABLE 2
Mechanical properties of white fibers for various stretch profiles.
				Young's
	Tow Denier	Tenacity	Elongation	Modulus
Sample ID	[g/9000 m]	[g/d]	[%]	[g/d]
WF-1	925	6.01	10.49	123.85
WF-2	919	5.51	11.57	98.46
WF-3	923	4.58	11.09	96.68

The linear densities of each WF sample were within 1% variability, which confirms that each given stretch profile produced the same diameter filaments (given the same precursor composition and bulk density). Interestingly, the white fiber tenacity and Young's Modulus increase in the order: WF-3<WF-2<WF-1. This behavior is consistent with results reported by Arbab et. al,⁶⁷ wherein higher 2^nd draw in relation to jet stretch was more effective at reducing porosity and increasing the stiffness of the fibers. Additionally, higher tenacity and modulus may be expected from the higher structural order and alignment of the polymer chains observed in WF-1.

Wet spinning of a PAN-based precursor was performed on a pilot spinning line with isolation of stretching variables from polymer composition, bath temperatures, coagulation conditions, and other process parameters. Three stretch distribution variants were selected with increasing jet stretch in exchange for a proportionate reduction in hot draw resulting in constant white fiber deniers and winder speeds. Three samples were collected for analysis for each stretch profile at the following locations: coagulation, first draw, and at the fiber winder. As determined by SEM, WAXS, DMA, and DSC, structural evolution occurs between sample locations, but final white fiber sample properties show key differences, including tenacities and moduli of the white fibers.

Increasing jet stretches result in higher shear forces and incomplete coagulation of the fibers, which leads to variable cross-sectional shapes and skin-core structures. In order of greatest orientation and crystallite thickness: Coag-1>Coag-2>Coag-3, which directly corresponds to longer residence times and lower shear forces in coagulation. This result propagates through to samples collected after first draw and samples collected at the winder with orientation increasing gradually between each stretching stage. The results show a strong correlation between orientation and the structural relaxation of the mesophase domain, which demonstrates the importance of complete coagulation and establishing ordered domains in the nascent stages of spinning. Further WF-1, with the greatest hot draw ratio, showed the most significant increase in L_c between the FD and WF stage, which corresponds well to the most substantial change in cyclization behavior. Cyclization activation energy was shown to correlate well with L_c, wherein higher ordered domains result in higher activation energy and delayed onset temperatures. The inventive process described herein employs the important findings that (1) coagulation influences initial structure and dictates orientation and structural evolution, (2) hot draw, or high temperature draw, post-wet stretch is beneficial for increasing the size of the mesophase domain, and (3) structural order directly impacts mechanical properties and downstream process parameters for TOS.

It would be apparent to a person of ordinary skill in that art that the conditions for conducting the inventive processes described herein may be optimized based on the intended application and circumstances without departing from the spirit of the present disclosure.

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Claims

1. A process for producing polymer fiber, the process comprising:

a) spinning a polymer solution into a coagulation bath, wherein a jet stretch is applied, to form coagulated fiber;

b) subjecting the coagulated fiber obtained in step (a) to a wet stretch to form a first drawn fiber; and

c) subjecting the first drawn fiber obtained in step (b) to a hot stretch, thereby forming the polymer fiber; wherein the amount of jet stretch, the amount of wet stretch, and the amount of hot stretch are effective to achieve the following properties wherein: the Herman's orientation factor of the polymer fiber produced is at least 0.60, and the crystallite thickness of the polymer fiber produced is at least 3 nm, greater than the crystallite thickness of the first drawn fiber.

2. The process according to claim 1, wherein the crystallinity of the coagulated fiber is no more than 8%, greater than the crystallinity of the first drawn fiber.

3. The process according to claim 1, wherein a linear mass density of the polymer fiber produced is from 0.7 to 1.2, denier per filament.

4. The process according to claim 1, wherein an average diameter of the coagulated fiber is at least 40 μm.

5. The process according to claim 1, wherein an average diameter of the first drawn fiber is at least 15 μm.

6. The process according claim 1, wherein the Herman's orientation factor of the coagulated fiber is at least 0.35.

7. The process according to claim 1, wherein the Herman's orientation factor of the polymer fiber produced is at least 0.08 greater than the Herman's orientation factor of the first drawn fiber.

8. The process according to claim 1, wherein an activation energy of structural relaxation for βc of the polymer fiber produced is less than 700 kJ/mol.

9. The process according to claim 8, wherein the activation energy of structural relaxation for βc of the polymer fiber produced is from 500 to 600 kJ/mol.

10. The process according to claim 1, wherein a cyclization activation energy of the polymer fiber produced is at least 7 kJ/mol greater than the cyclization activation energy of the first drawn fiber.

11. The process according to claim 1, wherein a tenacity of the polymer fiber produced is at least 4 g/d.

12. The process according to claim 1, wherein a Young's Modulus of the polymer fiber produced is at least 95 g/d.

13. The process according to claim 12, wherein the Young's Modulus of the polymer fiber produced is 95 to 130 g/d.

14. The process according to claim 1, wherein a peak temperature of a cyclization exotherm of the polymer fiber produced is at least 3° C. higher than the peak temperature of the cyclization exotherm of the coagulated fiber and/or the first drawn fiber.

15. The process according to claim 1, wherein the polymer fiber produced is polyacrylonitrile-based polymer fiber.

16. The process according to claim 1, wherein the spinning of the polymer solution is achieved by wet spinning.

17. The process according to claim 1, wherein the coagulation bath comprises a mixture of DMSO and water.

18. The process according to claim 1, wherein the polymer fiber produced is carbon fiber precursor fiber.

19. A polymer fiber produced according to the process of claim 1.

20. A process for producing carbon fiber, the process comprising oxidizing the polymer fiber according to claim 19 to form stabilized carbon fiber precursor fibers and then carbonizing the stabilized carbon fiber precursor fiber, thereby producing carbon fiber.