JP2024508516A - Process for making polymer fibers and polymer fibers made therefrom - Google Patents

Abstract

The present disclosure generally relates to processes for making polymeric fibers, typically polyacrylonitrile-based fibers, the characteristics of which include the amount of jet drawing, the amount of wet drawing, and the amount of hot drawing used. , controlled by certain parameters of the process. The present disclosure also relates to processes for making carbon fibers from such polymer fibers. [Selection diagram] None

Description

CROSS-REFERENCE TO RELATED APPLICATIONS 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.

The present disclosure generally relates to processes for making polymeric fibers, typically polyacrylonitrile-based fibers, the characteristics of which include the amount of jet drawing, the amount of wet drawing, and the amount of hot drawing used. , controlled by certain parameters of the process. The present disclosure also relates to processes for making carbon fibers from such polymer fibers.

Carbon fibers, particularly polyacrylonitrile (PAN)-based carbon fibers, are in increasing demand across many market sectors due to their high performance-to-weight ratio ^. However, their relative high cost has hindered market ubiquity and wide industry acceptance2,3 ^. To reduce costs, carbon fiber manufacturers are turning to larger tow sizes to increase production capacity, increase throughput, and reduce capital costs per pound of fiber processed. There are ^{4 to 7} . Although large tows (>24,000 filaments) have clear cost advantages, there are processing challenges, making wet spinning the preferred manufacturing method ^8,9 . Large tows result in increased filament breakage rates during hot drawing, ⁸ the inability to apply steam drawing or conventional winding mechanisms, ^9,10 higher heat generation during thermal oxidative stabilization (TOS), and carbon fibers. It suffers from further downstream challenges, including variable ^yields11 . Despite these additional challenges in processing large tows, the property goals and performance indicators have not decreased from their small tow counterparts.

To achieve carbon fiber property goals, emphasis has been placed on precursor development and fiber spinning to establish the initial carbon fiber structure and morphology ^{12 - 15} . The structure of the precursor is influenced by the chemical ^{composition16,17} , fiber denier and crystallinity18 ^, fibril structure and ^{orientation19-21} , and spinning ^speed20 . Structural defects caused by the spinning process, such as cavities, cracks, and scratches ^, can affect the shrinkage of the fibers in TOS and reduce the strength of the resulting carbon fibers. Therefore, elucidating the important properties of wet spinning in the development of PAN precursor structures can lead to improved mechanical properties of carbon fibers.

Although the morphology and structure of PAN are well defined in the literature ¹³ , researchers were initially perplexed by the differences observed in experiments under various process histories ^{25 - 28} . Bashir et al. carried out a series of impressive studies and revealed that the structure of atactic PAN can be divided into unoriented and oriented states ^{29 - 32} . In the unoriented state, PAN exhibits a two-phase morphology represented by an amorphous region and an ordered mesophase region, which is observed by two thermal transitions in differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA). Take ^{30, 32} . The lower temperature DMA transition (β _c ) of about 100 °C is thought to be attributed to the mesophase region, whereas the higher temperature transition (α) of about 130 to 160 °C is thought to be attributed to the amorphous region. . As the PAN fibers become oriented upon drawing and chain packing occurs, the intensity of the α peak decreases until it disappears, at which point the oriented PAN is ^{26, 31, 33, 34} .

In addition to thermal transitions and chain packing, the conformation of the chains has also been shown to shift when PAN transitions between unoriented and oriented states. Sawai et al. demonstrated that unoriented atactic PAN adopts a larger proportion of irregular helical arrangements, and upon stretching, the oriented chains rearrange into a planar zigzag structure ^. Shen et al. stated that the syndiotactic system with the highest degree of planar zigzag configuration has the lowest energy, highest stiffness, and the most favorable chain structure for intermolecular packing and TOS ^.

The structure of the precursor is clearly defined in the unoriented state and the oriented state, but the evolution of the microstructure of the PAN-based precursor or the path from the unoriented state to the oriented state in fiber spinning is directly linked. This has been little explored because it is capital intensive and difficult to mimic commercial spinning processes and investigate the effects of process changes on fiber structure. Rather, many studies have focused on the ^early stages of fiber spinning, where coagulation parameters can have a direct and significant impact on the precursor structure. In particular, much focus has been placed on the extrusion speed and jet stretching of the dope, or the stretching carried out between the spinneret and the first godet from coagulation in wet spinning ^{45 - 49} .

Apart from jet drawing, wet spinning includes other stages of drawing that can be used interchangeably to achieve a target filament denier before TOS and carbonization. Stretching can be applied immediately after solidification in a gel bath or some non-solvent medium, ^39,50,51 by conventional dry tensile stretching, ^37,52,53 or through a hot stretching process using hot liquid or vapor chambers ^. be able to. Each drawing step reduces the diameter of the fiber, but depending on the nature of the structure in the gel state versus the hot drawn state, the drawing mechanism can be very different. Edrington's thesis work demonstrated the limits of overall stretching in combinations of stretching steps, but separated jet stretching and gel stretching from hot stretching to determine which ones have the greatest impact on crystal alignment and orientation. However, it was not possible to identify the ^cause51 .

Therefore, there remains a need to achieve higher orientation states through optimal spinning processes that may result in more stable crystalline regions and precursor structures favorable for conversion into carbon fibers. A process for making polymeric fibers suitable for forming carbon fibers is described herein. The process of the present invention uses insights derived from wide-angle X-ray scattering (WAXS) and microstructural investigations of thermal transitions in DMA and stabilization kinetics by DSC of PAN-based polymer fibers.

Advantageously, it has been surprisingly discovered that the predetermined stages of stretching are not reciprocal and that the stretching profile can play an important role in the precursor structure. The process of the present invention uses insights derived from wide-angle X-ray scattering (WAXS) and microstructural investigation of thermal transitions in DMA and stabilization kinetics by DSC of polymeric fibers, such as PAN-based fibers.

In a first aspect, the present disclosure relates to a process for making polymeric fibers, the process comprising:
a) spinning a polymer solution into a coagulation bath to form coagulated fibers, in which jet drawing is applied;
b) wet-drawing the coagulated fibers obtained in step (a) to form first drawn fibers;
c) hot-stretching the first drawn fiber obtained in step (b), thereby forming a polymer fiber;
The amount of jet stretching, the amount of wet stretching, and the amount of hot stretching are effective to achieve the following properties:
The produced polymer fibers have a Herman orientation coefficient of at least 0.60, typically at least 0.65, more typically at least 0.67;
The crystallite thickness of the produced polymeric 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.

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

Figure 2 shows the structural evolution of fibers over the process described herein in which specific drawing profiles were used, as observed by WAXS analysis.

As used herein, the terms "a," "an," or "the" refer to "one or more," or "at least one," unless otherwise specified. ” and can be used interchangeably.

As used herein, the term "and/or" used in phrases of the form "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" is synonymous with "including," "containing," or "characterized by," and is intended to be inclusive or non-limiting. and does not exclude additional, unlisted elements or steps. The transition phrase "consisting essentially of" refers to materials or steps that, in addition to the specified materials or steps, essentially affect the essential characteristics or functions of the described composition, process, method, or product. It encompasses what is not given. The transition phrase "consisting of" excludes any element, step, or ingredient not specified.

Unless defined otherwise, 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. have

As used herein, 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; It depends in part on how it 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" refers to 50%, 20%, 15%, 10%, 9%, 8%, 7%, 6%, 5%, 4% of a given value or range. %, 3%, 2%, 1%, 0.5%, or within 0.05%.

It will also be understood that any numerical range recited herein is intended to include all subranges subsumed therein. For example, the range "1-10" is intended to include all subranges between and including the minimum recited value of 1 and the maximum recited value of 10, i.e. It has a minimum value of 1 or more and a maximum value of 10 or less. The numerical ranges disclosed are continuous and therefore include all values between the minimum and maximum values. Unless otherwise specified, the various numerical ranges specified in this application are approximations.

Throughout this disclosure, various publications may be incorporated by reference. To the extent that the meaning of any language in such publications incorporated by reference conflicts with the meaning of the language of this disclosure, the meaning of the language of this disclosure shall control, unless indicated otherwise.

A first aspect of the present disclosure relates to a process for making polymer fibers, the process comprising:
a) spinning a polymer solution into a coagulation bath to form coagulated fibers, in which jet drawing is applied;
b) wet drawing the coagulated fiber obtained in step (a) to form a first drawn fiber;
c) hot drawing the first drawn fiber obtained in step (b) to thereby form a polymer fiber;
The amount of jet stretching, the amount of wet stretching, and the amount of hot stretching are effective to achieve the following properties:
The produced polymer fibers have a Herman orientation coefficient of at least 0.60, typically at least 0.65, more typically at least 0.67;
The crystallite thickness of the produced polymeric 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.

In step a) of the process, a polymer solution, or "spinning dope", 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. Thus, in one embodiment, the polymer fibers made are polyacrylonitrile-based polymer fibers.

The polyacrylonitrile-based polymer can be any polymer containing repeating units derived from acrylonitrile. Suitable polyacrylonitrile-based polymers can be homopolymers consisting 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 commercial sources or prepared according to methods known to those skilled in the art. For example, the polymer can be produced by any polymerization method including, but not limited to, solution polymerization, dispersion polymerization, precipitation polymerization, suspension polymerization, emulsion polymerization, and variations thereof.

The polyacrylonitrile polymer comprises acrylonitrile and at least one comonomer selected from the group consisting of vinyl acids, vinyl esters, vinyl amides, vinyl halides, ammonium salts of vinyl compounds, sodium salts of sulfonic acids, and mixtures thereof. Contains repeating units derived from.

In one embodiment, the polyacrylonitrile-based polymer includes acrylonitrile, 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, vinyl imidazole ( VIM), acrylamide (AAm), diacetone acrylamide (DAAm), allyl chloride, vinyl bromide, vinyl chloride, vinylidene chloride, sodium vinylsulfonate, sodium p-styrenesulfonate (SSS), sodium methallylsulfonate (SMS) ), sodium 2-acrylamido-2-methylpropanesulfonate (SAMPS), and at least one comonomer selected from the group consisting of sodium 2-acrylamido-2-methylpropanesulfonate (SAMPS), and mixtures thereof.

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

The molar weight of polyacrylonitrile-based polymers suitable for use with the above process may be in the range 60-500 kg/mol, typically 90-250 kg/mol, more typically 115-180 kg/mol. .

Suitable solvents for the polymers are dimethyl sulfoxide (DMSO), dimethylformamide (DMF), dimethylacetamide (DMAc), ethylene carbonate (EC), N-methyl-2-pyrrolidone (NMP), zinc chloride (ZnCl ₂ ) /water, sodium thiocyanate (NaSCN)/water, and mixtures thereof, typically dimethyl sulfoxide (DMSO), dimethylformamide (DMF), dimethylacetamide (DMAc), ethylene carbonate (EC ), N-methyl-2-pyrrolidone (NMP).

The polymer solution used typically does not contain gels and/or aggregated polymers. The presence of gel and/or aggregated polymer can be determined using any method known to those skilled in the art. For example, a Hegman gauge can be used to determine the presence of gel and/or aggregated polymer. The polymer solutions used are generally stable and do not form gels over time.

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

The polymer solution can be subjected to conventional wet spinning and/or air gap spinning after removing air bubbles by vacuum. In wet spinning, the spinning dope is filtered and extruded through holes in a spinneret (typically metal) into a liquid coagulation bath in order for the polymer to form filaments. The spinneret holes determine the desired filament count of the fiber (eg, 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 coagulation bath. In one embodiment, spinning of the polymer solution is accomplished by wet spinning.

The coagulation liquid used in this process is a mixture of solvent and non-solvent. Water or alcohol is typically used as a non-solvent. Suitable solvents include those described herein. In one embodiment, dimethylsulfoxide, dimethylformamide, dimethylacetamide, or mixtures thereof are used as the solvent. In another embodiment, 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 coagulation rate of the extruded nascent filaments during coagulation. However, coagulation baths typically contain from 40% to 85% by weight of one or more solvents, with the remainder being non-solvent. In one embodiment, the coagulation bath includes a mixture of DMSO and water.

Jet stretching is applied in step a). As used herein, ^the amount of jet stretching applied in the coagulation step is the ratio of the first roller take-up speed to the dope extrusion speed, as understood by those skilled in the art. The dope extrusion rate _Vjet 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 , is the diameter of each hole. Accordingly, those skilled in the art will recognize that the V _-jet can be tuned by selecting appropriate values of Q, N, and D.

In step b) of this process, the coagulated fibers obtained in step (a) are subjected to wet drawing to form a first drawn fiber. As used herein, wet stretching is synonymous with "wet stretching," "first stretching," "first stretching," and "FD," and these terms are used interchangeably. can be done. During wet drawing, the coagulated fibers are immersed in one or more baths and are conveyed through the one or more baths, typically by the action of rollers or drawing rolls. As used herein, the amount of wet stretching is defined as the ratio of draw roll speeds between solidification and first stretching. Each of the baths or baths used for wet stretching may be maintained at a temperature of 40°C to 100°C.

In step c) of this process, the first drawn fiber obtained in step (b) is subjected to hot drawing, thereby forming a polymer fiber. As used herein, hot stretching is synonymous with "hot stretching," "second stretching," "second stretching," and "HD," and these terms are used interchangeably. obtain. During hot drawing, the first drawn fiber is conveyed through a heat source without being immersed in liquid, typically by the action of rollers or drawing rolls. For example, the heat source can be a steam furnace or a tube furnace through which the first drawn fiber obtained in step (b) is conveyed. As used herein, the amount of hot stretching is defined as the ratio of the stretch roll after FD to the final winding speed.

The processes of the present disclosure may be performed continuously or batchwise. As used herein, a process that is "continuously conducted" means that the fiber is passed through one or more processing steps at a time, without any interruption in time, material, or sequence. Refers to a process that carries a unit of work. This is in contrast to a batch process, in which a finite amount of material is processed at the end of a series of sequences that are carried out in a defined order and must be repeated to process or produce another batch of material. or as a process comprising one or more steps in a sequence of manufacturing steps. In one embodiment, the process is performed continuously.

Polymer fibers, typically polyacrylonitrile-based fibers, made by the processes described herein can be used as precursor fibers, so-called white fibers (“WF”), for the production of carbon fibers. Thus, in one embodiment, the polymer fibers made are carbon fiber precursor fibers.

In the process described herein, the amount of jet drawing, the amount of wet drawing, and the amount of hot drawing are effective to achieve the following properties, namely, the Herman orientation coefficient of the produced polymer fibers is , at least 0.60, typically at least 0.65, more typically at least 0.67, and the crystallite thickness of the produced polymer fiber is equal to the crystallite thickness of the first drawn fiber. at least 3 nm, typically at least 3.5 nm, more typically at least 4 nm.

Solidification influences the initial structure and dictates the orientation and structural evolution; hot stretching, or hot stretching, post-wet stretching is beneficial to increase the size of the mesophase region; and that structural regularity , has been found to directly influence the mechanical properties and process parameters downstream of TOS. Accordingly, the amount of jet stretching, the amount of wet stretching, and the amount of hot stretching are adjusted to achieve specific properties.

Wide-angle X-ray scattering (WAXS) is used to measure parameters used to determine various properties of fibers throughout the process, such as Hermann orientation coefficient, crystallite thickness, crystallinity, and D-spacing. .

WAXS measurements are performed according to methods known to those skilled in the art. In a case in point, WAXS is implemented in transmit mode using a Xenocs Xeuss 2.0 system. The light source is a GeniX3DCu ULD 8keV with a wavelength λ of 1.54189 Å. A single fiber tow (1,000 filaments) is aligned and secured across the aperture card. The aperture card with aligned fiber bundles is then transferred to the WAXS sample holder, which is placed 101.17 mm from the 2D detector (Dectris Pilatus 200K). The exposure time of the precursor fibers is 600 seconds. Data processing to obtain the integrated diffraction intensity for 2θ and azimuthal angle φ can be performed using the software Foxtrot provided by Xenocs. The structural parameters used were a Lorentzian function with two crystalline peaks located at approximately 2θ ₁ ≈16.9° and 2θ ₂ ≈29.3°, and an amorphous peak centered at approximately 2θ _a ≈25°. ^can be determined by peak fitting of Foxtrot data using Origin 2017 software.

The crystallinity χ _c is determined by Equation 2, where A _θ1 and A _θ2 are the areas of the two crystalline peaks 2θ ₁ and 2θ ₂ , respectively, and A _θa is the area of the amorphous peak. be. The D spacing is determined by Bragg's law (Equation 3), and the crystallite thickness L _c is determined by Scherrer's equation (Equation 4), where the shape factor K is 0.89 and β is the full width at half maximum in radians of the corresponding crystalline peak. The azimuthal scan profile is obtained by integrating the diffraction intensity I in the azimuthal direction with a (100) diffraction plane (2θ ₁ ≈16.9°) ^{55 , 56} . The Harman orientation coefficient (f _c ) ^is determined by equations 5 and 6.

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

In one embodiment, the produced polymer fibers have a Herman orientation coefficient of at least 0.60, typically at least 0.65, more typically at least 0.67.

In another embodiment, the coagulated fibers have a Herman orientation coefficient of at least 0.35, typically at least 0.40, more typically at least 0.42.

In yet another embodiment, the Hermann orientation coefficient of the polymer fibers made is at least 0.08, typically at least 0.1 greater than the Hermann orientation coefficient of the first drawn fiber.

In one embodiment, the crystallite thickness of the produced 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. .

In one embodiment, the crystallinity of the coagulated fibers is 8% or less, typically 7% or less, and more typically 6% or less than the crystallinity of the first drawn fiber.

Throughout the process, the fibers obtained at each step can be characterized by their cross-sectional diameter. Any suitable method of measuring fiber diameter can be used. In a suitable example, scanning electron microscopy (SEM) is used. In preparation for imaging, polymer fibers are lyophilized in a Labconco lyophilizer using a tert-butanol/water co-solvent system. The filament cross-section is produced by wetting the fiber tow with water and then immersing it in liquid nitrogen before breaking. The filament end is attached to a 15 mm aluminum SEM stub using carbon tape and sputter coated with 3 nm platinum to reduce charging. A Hitachi S-4800 can be used to image cross sections at 2 kV and various magnifications. Average filament diameter can be measured using QUARTZ PCI image processing software. At least 10 measurements are taken per sample.

In one embodiment, the average diameter of the coagulated fibers is at least 40 μm, typically at least 45 μm, more typically at least 50 μm, even more typically at least 55 μm.

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

The fibers obtained at each step of the process can also be characterized by the activation energy of structural relaxation ΔH at the mesophase glass transition temperature β _c .

The activation energy of structural relaxation ΔH at the mesophase glass transition temperature β _c is determined using dynamic mechanical analysis (DMA) using methods known to those skilled in the art. Suitably, DMA is performed on a TA Instruments Discovery Hybrid Series HR-2 rheometer. The precursor fiber tow is tested using a rectangular tension fixture configuration in tension mode. Condition the sample by applying a 5 minute temperature soak and an axial force of ±0.1N. Variables of loss modulus, storage modulus, and tan δ increase the oscillation temperature over a temperature range of 45 °C to 165 °C at 1 °C/min at various frequencies of 0.5, 1, 5, 10, and 15 Hz. plotted against temperature. Axial strain is set to 0.1%. The activation energy ΔH of structural relaxation at the mesophase glass transition temperature β _c obtained as the peak of tan δ is calculated by the Arrhenius relation with the frequency f and the general gas constant R according to Equation 7 ^. .

In one embodiment, the activation energy for structural relaxation of β _c of the prepared polymeric fibers 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 for βc structural relaxation of the prepared polymer fibers is between 500 and 600 kJ/mol, typically between 530 and 570 kJ/mol.

The fibers obtained at each step of the process can be characterized by their cyclization activation energy. 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 Instruments Q2000 with Universal Analysis 2000. The DSC was performed using certified reference materials (melting temperature: 156.60°C ± 0.03°C, enthalpy of fusion: 28.70 J/g ± 0.0°C, according to ASTM E967-03 and ASTM E968-02, respectively, for temperature and heat flow. 09 J/g) using indium metal. A standard aluminum DSC pan with a lid from TA instrument can be used with a nitrogen or air flow rate of 55 mL/min. The DSC is equilibrated at 35°C for 2 minutes and then ramp rates of 2, 5, 10, 20, and 30°C/min from 35 to 450°C are used.

The cyclization activation energy is determined from the Flynn-Wall-Ozawa (FWO) method, Equation 8, where E _a is the cyclization activation energy and R is the general gas where T _pk is the peak exothermic temperature (in Kelvin) and ω is the temperature ramp rate (in ^Kelvin ).

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

DSC can also be used to determine the peak temperature of the cyclization exotherm of the fibers obtained at each step of the process. The cyclization reaction is generally exothermic and can be seen as a peak in a DSC scan where the heat flux (in W/g) is plotted as a function of temperature. In one embodiment, the peak temperature of the cyclization exotherm of the produced polymeric fibers is at least 3° C. higher than the peak cyclization exotherm temperature of the coagulated fibers and/or the first drawn fibers.

The polymeric fibers formed by the inventive processes described herein, typically carbon fiber precursor fibers, or white fibers, have different draw profiles, i.e., the amount of jet drawing, the amount of wet drawing, and the amount of hot drawing. derived from the surprising observation that the amount of carbon gives rise to specific microstructures with mechanical properties such as toughness, elongation, and Young's modulus.

The linear mass density, or tow denier, of the formed polymer fibers can be calculated from the average of two yield measurements based on the weight of a 3 meter sample section. Linear mass density can be expressed on a tow (bundle of filament) basis or on a per filament basis. Linear mass density is expressed herein on a filament-by-filament basis. Therefore, linear mass density as used herein is mass in grams per 9000 meters of filament, or denier per filament. The linear mass density of the polymer fibers produced is between 0.7 and 1.2 denier per filament, typically between 0.85 and 1.0 denier.

Mechanical properties can be tested using known methods. For example, MTS Criterion C43 and Testworks 4 software can be used. The 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 three times at a crosshead speed of 2 inches/minute.

In one embodiment, the toughness of the polymer fibers made is at least 4 g/d, typically at least 5 g/d, more typically at least 6 g/d.

In one embodiment, the Young's modulus of the polymer fibers made is at least 95 g/d, more typically at least 100 g/d.

In another embodiment, the Young's modulus of the polymer fibers made is between 95 and 130 g/d, typically between 100 and 130 g/d, and more typically between 115 and 125 g/d.

In a second aspect, the present disclosure relates to a process for making carbon fibers, the process oxidizing a polymer fiber described herein or a polymer fiber made according to a process described herein. and forming stabilized carbon fiber precursor fibers, and then carbonizing the stabilized carbon fiber precursor fibers to thereby produce carbon fibers.

Polymer fibers formed according to the processes described herein can be oxidized to form stabilized carbon fiber precursor fibers, and then the stabilized carbon fiber precursor fibers are carbonized, Produce carbon fiber.

During the oxidation stage, the polymer fibers are passed under tension through one or more specialized ovens, each having a temperature of 150-300°C, typically 200-280°C, where heated air is supplied to each oven. The oxidation process causes oxygen molecules in the air to bond with the fibers and initiate cross-linking of the polymer chains, thereby increasing the fiber density to the appropriate level. Such oxidized fibers, typically PAN fibers, have an insoluble ladder aromatic molecular structure and are ready for carbonization processing.

Carbonization results in crystallization of carbon molecules, resulting in finished carbon fibers having a carbon content of greater than 90 percent. Carbonization of the oxidized or stabilized carbon fiber precursor fibers is carried out in one or more specially designed furnaces in an inert (oxygen-free) atmosphere, typically under a nitrogen atmosphere. The oxidized carbon fiber precursor fibers are passed through one or more ovens, each heated to a temperature between 300°C and 1650°C.

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

The carbon fibers can then be subjected to a sizing process in which a size coating (eg, an epoxy-based coating) is applied to the fibers. Sizing can be carried out by passing the fiber through a sizing bath containing a liquid coating material. The sizing process protects the carbon fibers during handling and processing into intermediate forms, such as dry fabrics and prepregs. The sizing process also aggregates the filaments into individual tows to reduce fuzz, improve processability, and increase interfacial shear strength between the fibers and matrix resin used to make the composite. After the sizing process, the coated carbon fibers are dried and then wound onto a bobbin. The produced carbon fibers are suitable for use in making composite materials.

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

Example 1. Preparation of Polymer Fibers A polymer dope was prepared by mixing the PAN-based polymer in DMSO in a 15 gallon Myers mixer to a target concentration of 18.5% solids by weight. The polymer dope was warmed 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 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 (hole diameter = 50 μm) into a heated coagulation bath containing a mixture of DMSO and water.

Polymer composition, bath concentration, temperature, and other spinning parameters remained constant; the only variable was the draw ratio performed during the jet draw and the second draw. As outlined in Table 1 below, three variants were considered with proportional amounts of stretching performed between the jet stretch and the second stretch, where the variants , "1", "2", or "3", and the process at which the fiber sample was taken is "Coag" (after coagulation), "FD" (after first drawing), and "WF" ( after the second stretching). The amount of jet stretching applied during solidification was controlled by the ratio X of the first roller take-up speed to the dope extrusion speed ^. The dope extrusion speed _Vjet is calculated by Equation 1 above. The ratio of draw roll speeds between coagulation ("Coag") and first draw ("FD") is the ratio between the amount of wet drawing, Y, and the draw rolls after FD, and the final winder Speed is the amount of stretching Z applied in hot stretching.

For Coag-2 and Coag-3, jet draw increased by 36% and 58%, respectively, and second draw decreased by a similar amount compared to Coag-1. For all three variants considered, the final winder speeds were similar, so the total stretch applied to each white fiber precursor was similar X*Y*Z. For each of the three variants, samples collected after solidification and first stretching were hung to dry to constant weight before further analysis. White fiber (WF) samples were collected in a winder after the drying process during spinning and used as is.

Example 2. SEM Analysis SEM analysis was performed on the fiber samples shown in Table 1 as described above.

It was observed that as the jet drawing increased from Coag-1 to Coag-2 to Coag-3, the cross-section of the filament transitioned from kidney bean shape to elliptical circular shape and finally to a more circular shape. . Jet stretching is known to affect shear forces in the spinneret capillary, which reduces die swelling and, as a result, solvent exchange ^. Also, as jet drawing increases, the diameter of the filaments exiting the spinneret capillary decreases, and the larger the protofiber radius, the thinner and sparser the film structure becomes, facilitating the interdiffusion of water and DMSO ^. At larger diameters, Coag-1 has a more gradual transition in the skin core and more distinct skin layers compared to Coag-2 and Coag-3. In Coag-3, compared to Coag-1 and Coag-2, interdiffusion is sufficiently inhibited to form a solid film, preventing complete solidification of the core structure and leaving a loose spongy character. In addition, higher jet stretching may cause disruption due to shorter residence time in the coagulation bath, which may leave the fibers not fully coagulated.

After solidification, the fibers were all subjected to the same draw ratio in the first draw, although the actual speed and residence time varied due to the different jet draws. Although the cross-sectional shape converges towards circular filaments in all three cases, FD-1 retains some kidney bean similarity. Cross-sectional images are obtained by fracturing samples in liquid nitrogen and can provide insight into defects from the fractured surface ⁶¹ . FD-1 and FD-2 show a radiation pattern from the center of the fiber, which is indicative of internal defects. FD-3 shows separation of the film layer from the core, suggesting that the fracture started at the film and spread around the core of the filament rather than through the center of the filament. ing. Also, the film of the filament shows obvious differences, with FD-1 having finer stripes along the fiber axis, whereas FD-3 has more pronounced and deeper stripes due to faster solidification rate. ⁶² with membrane grooves. The average filament diameter maintains the same proportional difference after the first draw as in solidification, as expected at a constant first draw ratio. For example, the average filament diameters of Coag-2 and FD-2 (49.6 μm and 19.1 μm) remain approximately 86% of Coag-1 and FD-1 (58.0 μm and 22.2 μm), respectively. A proportional adjustment in the second draw is used to compensate for speed differences so that the winder speeds for each of the three profiles are equal. The average diameters of the coagulated fibers and first drawn fibers as observed by SEM are summarized in Table 2 below. Morphological examination of the WF samples by SEM is difficult due to the ductile nature of the fibers and is therefore not included.

Example 3. WAXS Analysis WAXS was performed on each sample to investigate the differences in the structure of the fibers after each stage of drawing. Figure 1 shows the structural evolution of Profile 1, with the scans of each sampling location superimposed. Interestingly, Coag-1 exhibits a strong peak around 2θ ₁ ≈17°, which corresponds to the mesophase regularity region in PAN, ³¹ indicating that a significant amount of crystallization occurs as a result of solidification and jet stretching. It is shown that. Coag-1 also exhibits a broad amorphous peak at about 2θ _a ≈25° and the ^onset of another more highly ordered diffraction plane at about 2θ ₂ ≈29°. When the fiber is drawn further in the first draw, FD-1 shows a slight narrowing of the 2θ ₁ and greater intensity of the 2θ ₂ peak compared to Coag-1. Following the second stretch, WF-1 exhibits a sharper 2θ ₁ diffraction peak, a very prominent 2θ ₂ peak, and a significant decrease in relative 2θ _a intensity.

Crystallinity (χ _c ), crystallite thickness (L _c ), and orientation (f _c ) under all processing conditions were determined as described above. Crystallinity appears to decrease throughout the process, but this behavior is believed to be an anomaly due to the data processing techniques used. Although the interpretation of crystallinity is often debated in the literature, ¹³ the proposed method used herein can be observed in the WAXS peak parameters and peak fits of various fiber samples (not shown). (b) Both the 2θ ₁ peak and the 2θ ₂ peak are considered as part of the regularity region. Nevertheless, in the amorphous region, the apparent crystallinity upon solidification is high due to the difficult deconvolution of _2θ2 by _2θa , and in samples with high amorphous content, the peak area at the _2θ2 peak is , which means that the overall crystallinity of the sample may be intentionally increased.

Surprisingly, the prominent 2θ ₁ peak of the coagulated sample indicates that the mesophase region is established almost at the beginning of spinning. Although the 2θ ₁ (d _2θ1 ) crystallinity and d-spacing remain relatively stable throughout the drawing process during spinning, large structural changes occur in crystallite thickness and orientation. For all three samples, the orientation increases gradually from Coag to FD to WF, with the peak narrowing with each azimuthal scan. The crystallite thickness has a slight increase from Coag to FD and a sharp increase from FD to WF noted by the narrowing of the _2θ1 peak. This suggests that hot stretching is more effective than wet stretching in fusing regular regions together and assimilating amorphous and regular regions. Liu et al. suggested that the early stages of stretching are devoted to straightening and disentangling the polymer chains, ²² which may also explain the marginal increase in L _c during wet stretching. The WAXS parameters of the fiber samples are summarized in Table 3 below.

Example 4. Molecular dynamics simulations Molecular dynamics simulations were performed to visualize the evolving structure suggested by the WAXS data. Here, a PAN system consisting of 20 individual chains was subjected to axial strain to simulate the stretching process.

Schrodinger Materials Science 19-3 and OPLS3-e force field were used for all simulations. The PAN chain was prepared using the "Crosslink Polymer" module by first placing 5000 acrylonitrile monomers and 20 initiator molecules into a periodic box, followed by a subsequent 50 ns NPT and NVT at 400 K. Constructed by performing stages to densify and relax cells. Dummy atoms were used to represent propagating radicals. Corresponding to the number of initiators, 20 reactions were carried out per step, but the monomer had to be within a distance criterion of 4.0 A to the dummy atom for the reaction to take place. After each step, the system was relaxed for 50 ps. Polymerization was carried out at 400 K and continued until the monomer conversion was equal to 100%. Following polymerization, the system was annealed at 400 K for an additional 50 ns of NPT and NVT. Stretching simulations were performed by expanding one dimension of the periodic box at 0.1% strain in 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 periodic box. Stretching was performed under NVT conditions. The torsion of the framework was analyzed according to a stretching simulation in which the torsion was defined by the dihedral angles of four atoms from four connected framework atoms.

A snapshot of the simulated system was observed in which a single chain was represented by an atom augmented to track its progress. The 0% system was omitted as it was completely amorphous and was not considered to adequately represent the actual polymer. The enlarged region of the 10% and 60% strain systems shows the local regularity of the atatic polymer chains, where at lower strains more population of gauche structures are present within the chains. It acts as a twist and the local regularity is disturbed. Observation of the simulation results highlighted the skeletal dihedral angles of various populations, but two configurations, trans (≈180°) and gauche (≈60°), were largely prioritized. As the stretching progresses, the gauche structure transforms into a trans structure, resulting in an expansion of the zigzag arrangement in the plane, allowing adjacent chains to align and condense, thus forming a larger mesophase region. . With further stretching, the gauche structure tends to zero and the system approaches single-phase behavior.

Example 5. DMA Analysis of Fibers The storage modulus (E′) and loss modulus (E”) and the corresponding tan delta peak evolution in the 10 Hz frequency scan of profile 1 were investigated using DMA analysis as described above, respectively. Because it is difficult to accurately measure the cross-sectional area of various samples, the magnitude of the storage modulus and loss modulus should not be considered important. Rather, the magnitude of the storage modulus and loss modulus should not be considered important. A tan delta is used for sample comparisons to eliminate sample size bias.

Coag-1 exhibited two peaks in the tan-delta curve, an α transition peak at about 145°C, and a β _c transition peak at about 105°C ^. When the fiber is oriented in FD-1, the α transition disappears and the magnitude of the β _c transition initially increases, with the α peak corresponding to the amorphous region and β _c corresponding to the ordered mesophase region. ³¹ . Coag-1 also exhibits a narrower β _c transition peak compared to FD-1 due to the presence of larger amorphous regions that are more pronounced in the as-spun fibers ⁶³ . As stretching progresses through WF-1, only a single β _c peak is present, further supporting the transition to single-phase behavior.

The activation energy of the β _c transition was determined by an Arrhenius fit that plots the inverse of the peak temperature 1/β _c as a function of the logarithm of the frequency ln f. The results are summarized in Table 4 below.

As the frequency increases, the transition temperature increases and the magnitude of tan δ decreases due to the shortened response time and the viscoelastic nature of PAN-based fibers. The magnitude of the slope of drawing profile 1 follows in the order Coag-1>FD-1>WF-1, which corresponds to a decreasing value of ΔH as the fiber is drawn.

Example 6. DSC Analysis DSC analysis was performed on fiber samples obtained from the process described herein. DSC results for fiber evolution at 20° C./min in nitrogen were determined. A clear onset peak associated with the exothermic cyclization reaction is observed at about 243° C. for Coag-1, which decreases for FD-1 and is absent for WF-1. This early peak is believed to be due to the larger size amorphous regions of Coag-1 and FD-1 compared to WF-1. WF-1 shows a slower onset of exotherm, which is probably due to the ordered regions and non-regular regions compared to Coag-1 and FD-1, which may have more separated phase boundaries. This is because the mixing of crystalline regions is greater. Due to the occurrence of a larger onset peak below 250 °C and a higher internal exotherm, the cyclization exotherm in the regular region is also lower than that of Coag-1 and FD-1, about 296 °C, compared to WF-1, about 299 °C. peaks at lower temperatures and higher heat flows in °C.

The activation energy of cyclization E _a is determined from the Ozawa method in which the reciprocal of the peak exothermic temperature 1/T _pk is plotted against the logarithm of the heating rate ln ω, increasing at faster ramp rates. published along with ^the peak temperature59,64. The peak temperatures and calculated E _a values are summarized in Table 5 below.

The cyclization activation energy remains approximately constant between the sample collected after solidification and the sample after the first stretch, with values ranging from 155 kJ/mol to 158 kJ/mol. The most significant change occurs for WF-1 during hot stretching, where the E _a value increases to 169.8 kJ/mol. The initiation of cyclization is believed to be due to structural anomalies and defects, ^64-66 which may explain why WF samples have much higher activation energies. The white fiber after hot drawing contains regular regions of maximum size, which leads to fewer structural defects at the onset of cyclization, as evidenced by WAXS and DMA analysis.

Example 7. Mechanical Properties of White Fibers Mechanical properties were used for the final fiber comparisons and are summarized in Table 6 below.

The variation in linear density for each WF sample was within 1%, confirming that each given drawing profile produced filaments of the same diameter (given the same precursor composition and bulk density). ). Interestingly, the toughness and Young's modulus of white fibers increase in the following order: WF-3<WF-2<WF-1. This behavior is consistent with the results reported by Arbab et ^al.67 , where a higher second draw with respect to jet drawing was more effective in reducing fiber porosity and increasing stiffness. . In addition, higher toughness and elastic modulus can be expected from the higher structural regularity and alignment of polymer chains observed in WF-1.

Wet spinning of PAN-based precursors was performed in a pilot spinning line, separating drawing variables from polymer composition, bath temperature, coagulation conditions, and other process parameters. Instead of proportionally decreasing hot drawing, we selected three draw classification variants that increased jet drawing to keep the white fiber denier and winder speed constant. For analysis of each draw profile, three samples were collected at the following locations: coagulation, first draw, and fiber winder. Although structural evolution occurs between sample locations, as determined by SEM, WAXS, DMA, and DSC, the properties of the final white fiber sample are affected by important factors such as white fiber toughness and elastic modulus. shows the difference.

Increased jet stretching increases shear forces and causes incomplete coagulation of the fibers, which changes the cross-sectional shape and film core structure. The order of orientation and crystallite thickness is as follows: Coag-1>Coag-2>Coag-3. This directly corresponds to longer residence times and lower shear forces in solidification. This result is transmitted to samples collected after the first draw and in the winder where the orientation gradually increases between each draw stage. The results show a strong correlation between the orientation of the mesophase region and structural relaxation, demonstrating the importance of complete solidification and establishment of ordered regions in the early stages of spinning. Furthermore, WF-1 with the highest hot drawing ratio showed the most significant increase in _Lc between the FD and WF stages, which corresponds well to the most substantial change in cyclization behavior. ing. The cyclization activation energy was shown to correlate well with L _c , while regions of higher regularity result in higher activation energies and delayed onset temperatures. The inventive process described herein (1) solidification affects the initial structure and dictates orientation and structure evolution; (2) hot stretching, or hot stretching, post-wet stretching controls the size of the mesophase region; (3) Structural regularity directly influences TOS mechanical properties and downstream process parameters.

It will be apparent to those skilled in the art that the conditions for implementing the inventive methods described herein can be optimized based on the intended application and environment without departing from the spirit of the disclosure.

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Claims (20)

a) spinning a polymer solution into a coagulation bath to form coagulated fibers, in which jet drawing is applied;
b) wet drawing the coagulated fiber obtained in step (a) to form a first drawn fiber;
c) A method for producing a polymer fiber, comprising a step of hot-drawing the first drawn fiber obtained in step (b) to thereby form a polymer fiber,
The amount of jet stretching, the amount of wet stretching, and the amount of hot stretching are determined by the following characteristics:
The polymer fibers produced have a Herman orientation coefficient of at least 0.60, typically at least 0.65, more typically at least 0.67;
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. methods that are effective in achieving this. 2. The crystallinity of the coagulated fiber is 8% or less, typically 7% or less, more typically 6% or less than the crystallinity of the first drawn fiber. Method. A method according to claim 1 or 2, wherein the linear mass density of the polymer fibers produced is between 0.7 and 1.2 denier, typically between 0.85 and 1.0 denier per filament. 4. The average diameter of the coagulated fibers is at least 40 μm, typically at least 45 μm, more typically at least 50 μm, even more typically at least 55 μm. Method. A method according to any preceding claim, wherein the average diameter of the first drawn fibers is at least 15 μm, typically at least 20 μm, typically at least 22 μm. 6. A method according to any preceding claim, wherein the Harman orientation coefficient of the coagulated fibers is at least 0.35, typically at least 0.40, more typically at least 0.42. 7. The Hermann orientation coefficient of the polymer fibers produced is at least 0.08, typically at least 0.1 greater than the Hermann orientation coefficient of the first drawn fibers. Method described. 8. The activation energy of β _c structural relaxation of the polymer fibers produced is less than 700 kJ/mol, typically less than 650 kJ/mol, more typically less than 600 kJ/mol. The method described in any one of the above. The method according to any one of claims 1 to 8, wherein the activation energy for structural relaxation of β _c of the polymer fibers produced is between 500 and 600 kJ/mol, typically between 530 and 570 kJ/mol. . The cyclization activation energy of the polymer fibers 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. A method according to any one of claims 1 to 9, wherein the method is molar large. A method according to any one of claims 1 to 10, wherein the toughness of the polymer fibers produced is at least 4 g/d, typically at least 5 g/d, more typically at least 6 g/d. . A method according to any preceding claim, wherein the Young's modulus of the polymer fibers produced is at least 95 g/d, more typically at least 100 g/d. Any one of claims 1 to 12, wherein the Young's modulus of the polymer fibers produced is between 95 and 130 g/d, typically between 100 and 130 g/d, more typically between 115 and 125 g/d. The method described in. Any one of claims 1 to 13, wherein the peak temperature of the cyclization exotherm of the produced polymer fiber 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 method described in. The method according to any one of claims 1 to 14, wherein the polymer fibers produced are polyacrylonitrile polymer fibers. A method according to any one of claims 1 to 15, wherein spinning of the polymer solution is achieved by wet spinning. 17. A method according to any one of claims 1 to 16, wherein the coagulation bath comprises a mixture of DMSO and water. A method according to any one of claims 1 to 17, wherein the polymer fibers produced are carbon fiber precursor fibers. Polymer fibers made according to the method according to any one of claims 1 to 18. oxidizing a polymeric fiber according to claim 19 or a polymeric fiber made according to the method according to any one of claims 1 to 18 to form a stabilized carbon fiber precursor fiber; and then carbonizing the stabilized carbon fiber precursor fibers to thereby produce carbon fibers.

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