High-strength and high-plasticity high-carbon steel and method for manufacturing the same

Abstract

The application belongs to the technical field of metal material heat treatment, and particularly relates to a high-strength and high-plasticity high-carbon steel and a preparation method thereof, the preparation method comprising the following steps: S1, obtaining high-carbon steel, and obtaining austenitic high-carbon steel by austenitizing the high-carbon steel; S2, obtaining martensitic high-carbon steel by rapidly cooling the austenitic high-carbon steel to 50-150 DEG C, wherein the cooling rate of the rapid cooling is greater than or equal to 50 DEG C / s, and the structure of the martensitic high-carbon steel comprises lath martensite; and S3, obtaining high-strength and high-plasticity high-carbon steel by tempering the martensitic high-carbon steel, wherein the tempering temperature is 320-380 DEG C, and the structure of the high-strength and high-plasticity high-carbon steel comprises nanolayered structure formed by alternately arranging ferrite lamellas and carbide nanolamellar layers. The high-strength and high-plasticity high-carbon steel composed of the nanolayered structure formed by alternately arranging ferrite lamellas and carbide nanolamellar layers is prepared by synergistically controlling the composition and process of the high-carbon steel, and based on the 'precursor preparation + low-temperature assembly' path of accurate control of phase transition kinetics.

Description

Technical Field

[0001] This application relates to the field of heat treatment technology for metallic materials, specifically, to a high-strength, high-ductility high-carbon steel and its preparation method. Background Technology

[0002] High-carbon steel (carbon content approximately 0.75%-0.85%) wire rod is a key raw material for manufacturing high-performance prestressed steel strands, wire ropes, and other wire products. Since the application of Stellmor controlled cooling technology in the last century, its standard production process has been basically finalized: after hot rolling, it is cooled at a controlled rate (usually 1-10℃ / s) on the Stellmor air-cooling line, with the goal of obtaining a refined sorbite (a refined pearlite) structure through a diffusion-type phase transformation from austenite to pearlite.

[0003] Although the process is mature and stable, its inherent phase transition kinetics bottleneck has led to the material properties approaching the ceiling of traditional approaches, specifically manifested in: 1. Limits to Microstructure Refinement and Performance Bottlenecks: The interlamellar spacing (λ) of pearlite is determined by the phase transformation undercooling (ΔT), following the Zener-Hillert relationship (λ∝1 / ΔT). Under current industrial conditions, the undercooling provided by air cooling is limited. Even with strong air cooling, the interlamellar spacing is typically limited to above 100 nanometers (mostly in the 120-200 nanometer range) because the cooling curve cannot avoid the pearlite transformation "nose" region (approximately 550-700℃). This relatively coarse microstructure leads to an irreconcilable contradiction between strength (tensile strength approximately 1100-1200 MPa) and plasticity (reduction of area typically <40%), especially during subsequent deep cold drawing processes, where coarse, brittle cementite lamellae easily become crack initiation sites, resulting in a persistently high wire breakage rate.

[0004] 2. Sensitivity to Process Control and Performance Fluctuations: The final microstructure of traditional processes is highly dependent on the cooling rate and temperature uniformity. Differences in cooling along the cross-section and length of the wire rod can easily lead to inconsistent core and surface microstructures (such as the appearance of troostite or even bainite in the core), and performance fluctuations at both ends, severely affecting batch-to-batch stability. Furthermore, this process is extremely sensitive to fluctuations in steel composition (especially carbon and manganese content); even slight compositional shifts can cause significant changes in phase transformation points and microstructure.

[0005] 3. Limitations of Existing Improvement Technologies: To overcome the aforementioned bottlenecks, the industry has attempted various improvement methods, but their fundamental approach remains confined to promoting the direct transformation of austenite to pearlite. For example, microalloying (adding V, Nb, etc.) refines the original grains by pinning austenite grain boundaries, indirectly refining pearlite clusters, but its direct impact on lamellar spacing is limited. Optimizing cooling parameters (such as lowering the spinning temperature and increasing the airflow) can increase supercooling to some extent, but its increase has already reached the boundaries of equipment capacity and phase transformation thermodynamics. Furthermore, further acceleration can easily lead to the bainite transformation region, forming a mixed structure and deteriorating overall performance. Essentially, these technologies have not broken away from the traditional paradigm of "high-temperature austenitization → medium-temperature pearlite phase transformation," and have failed to fundamentally innovate the phase transformation mechanism.

[0006] 4. Shortcomings of Existing Martensitic Tempering Processes: Tempering after quenching to obtain martensite (i.e., tempering) is a conventional method for achieving a good balance of strength and toughness in low and medium carbon steels. However, when this approach is applied to high carbon steels such as 82B, traditional high-temperature tempering (typically >450℃) often leads to significant spheroidization and coarsening of carbides, forming a uniformly distributed granular carbide structure (tempered sorbite). While this structure can improve toughness, it is difficult to retain and refine lamellar carbides, especially failing to obtain a long-range ordered, uniformly oriented nanoscale layered structure. Low-temperature tempering (<300℃), on the other hand, suffers from insufficient carbon atom diffusion, making it difficult to achieve sufficient precipitation and ordered arrangement of carbides. Therefore, existing technologies lack a systematic method for obtaining a microstructure in high carbon steel that combines a nanoscale lamellar structure with excellent strength-ductility balance through precise control of the tempering process.

[0007] Therefore, the current technological gap lies in the need for a completely new process route that can: ① completely avoid the traditional pearlite diffusion phase transformation, fundamentally breaking through the kinetic limitations of lamellar refinement; ② achieve nanoscale lamellar self-assembly of carbides during martensitic tempering through a specific combination of process parameters, rather than the traditional spheroidization or random precipitation; ③ obtain a unique microstructure that has a layered morphology similar to pearlite, but is significantly superior to traditional pearlite and conventional tempered structures in terms of formation mechanism, lamellar refinement, and performance. Summary of the Invention

[0008] To solve the above technical problems, the present application provides a method for preparing high-strength and high-plasticity high-carbon steel, comprising the following steps: S1. Obtain high-carbon steel, and subject the high-carbon steel to austenitization to obtain austenitic high-carbon steel. By mass percentage, the composition of the high-carbon steel is: C: 0.75% - 0.85%, Mn: 0.5% - 1.0%, Cr: 0.1% - 0.5%, Si: 0.1% - 0.5%, and the balance is Fe and inevitable impurities; S2. Rapidly cool the austenitic high-carbon steel to 50 - 150°C to obtain martensitic high-carbon steel. Among them, the cooling rate of the rapid cooling is ≥50°C / s, and the structure of the martensitic high-carbon steel includes lath martensite, and the content of the lath martensite is ≥90wt%; S3. Temper the martensitic high-carbon steel to obtain high-strength and high-plasticity high-carbon steel. The tempering temperature is 320 - 380°C, and the structure of the high-strength and high-plasticity high-carbon steel includes a nano-layered structure in which ferrite lamellae and carbide nano-lamellae are alternately arranged.

[0009] As a preferred scheme of the method for preparing high-strength and high-plasticity high-carbon steel described in the present application, in step S1, the holding temperature of the austenitization is 1000 - 1050°C, and the holding time of the austenitization is 20 - 40 min.

[0010] As a preferred scheme of the method for preparing high-strength and high-plasticity high-carbon steel described in the present application, in step S2, the cooling rate is determined according to the equivalent diameter D of the workpiece: when D ≤ 10 mm, the cooling rate is ≥50°C / s; when 10 mm < D ≤ 20 mm, the cooling rate is ≥80°C / s; when D > 20 mm, the cooling rate is ≥100°C / s.

[0011] As a preferred scheme of the method for preparing high-strength and high-plasticity high-carbon steel described in the present application, in step S3, when the tempering temperature is ≥320°C and ≤350°C, the tempering time is ≥5 h and ≤8 h.

[0012] As a preferred scheme of the method for preparing high-strength and high-plasticity high-carbon steel described in the present application, in step S3, when the tempering temperature is > 350°C and ≤380°C, the tempering time is ≥3 h and < 5 h.

[0013] [[ID=​​​​​This application also provides a high-strength, high-ductility, high-carbon steel, which is prepared using the above-described method for preparing high-strength, high-ductility, high-carbon steel.

[0016] As a preferred embodiment of the high-strength and high-ductility high-carbon steel described in this application, the microstructure of the high-strength and high-ductility high-carbon steel includes a nanolayered structure composed of alternating ferrite lamellae and carbide nanolamellae, wherein the average interlamellar spacing of the nanolayered structure is ≤80nm, and the aspect ratio of the carbide nanolamellae in the nanolayered structure is ≥5:1.

[0017] As a preferred embodiment of the high-strength, high-ductility, high-carbon steel described in this application, the high-strength, high-ductility, high-carbon steel has a tensile strength ≥1300MPa, a yield strength ≥1100MPa, a reduction of area ≥50%, and a uniform elongation ≥8.0%.

[0018] The beneficial effects of this application are as follows: This application provides a method for preparing high-strength and high-ductility high-carbon steel. By synergistically controlling the composition and process of high-carbon steel, and based on the precise control of phase transformation kinetics through the "precursor preparation + low-temperature assembly" path, high-strength and high-ductility high-carbon steel composed of a nanolayered structure with alternating ferrite sheets and carbide nanosheets is prepared.

[0019] This application employs compositional design to ensure a sufficiently high carbon content to form a supersaturated solid solution, while using lower alloying element contents to avoid excessively affecting the phase transformation point, thereby achieving a wider controllable process window. Sufficient C ensures adequate carbon content, providing the material basis for subsequent nano-carbide precipitation; moderate Mn improves hardenability, ensuring martensite rather than pearlite is obtained during rapid cooling; trace amounts of Cr enhance carbide stability and delay premature spheroidization; appropriate Si inhibits cementite coarsening, helping to maintain the lamellar structure; the synergistic process control of C, Mn, Cr, and Si, with cooling at a rate not less than 50℃ / s, ensures that the cooling curve completely avoids the pearlite-bainite transformation region, directly entering the martensite transformation region, thus obtaining a homogeneous, diffusionless single lath martensite precursor. Using 320-380℃ as the tempering temperature, within this range, carbon atoms possess moderate diffusion capabilities, enabling preferential growth of carbides along specific directions after nucleation, forming continuous and straight nanosheets while effectively suppressing their spheroidization tendency. Below 320℃, carbon atom diffusion capabilities are insufficient, making it difficult to achieve sufficient precipitation and long-range ordered arrangement of carbides, resulting in incomplete microstructure transformation and difficulty in forming nanolayered structures. Above 380℃, carbides will undergo significant spheroidization and coarsening, disrupting the continuity of the lamellar structure and causing the microstructure to degenerate into traditional granular tempered sorbite, losing its nanolayered characteristics.

[0020] The scheme of adjusting the cooling rate according to different workpiece diameters is to ensure that the core of large-sized workpieces can also obtain sufficient cooling rate, avoid incomplete transformation of the core structure, and fully consider the coupling of size effect and phase transformation kinetics.

[0021] This application guides carbides to "self-assemble" and grow along specific crystallographic directions by precisely controlling the tempering temperature and time, forming a long-range ordered nanosheet structure. This results in continuous, straight nanoscale sheet-like carbides (aspect ratio ≥ 5:1), with a morphology closer to pearlite but with a significantly finer interlamellar spacing. While maintaining ultra-high strength (≥ 1300 MPa), it significantly improves plasticity (reduction of area ≥ 50%), breaking through the bottleneck of the traditional quenching and tempering process where strength and plasticity are inversely related. Traditional quenching and tempering (quenching + high-temperature tempering) aims to spheroidize and homogenize carbides to obtain granular tempered sorbite, thus improving toughness, but this results in significant strength loss and fails to form sheet-like carbides.

[0022] The microstructure obtained in this application is the product of decomposition of a single supersaturated phase at intermediate temperature (α′→α+Fe3C, carbides precipitate from the supersaturated solid solution). The interlamellar spacing is precisely controlled by the carbide precipitation kinetics, breaking through the traditional refinement limit. By precisely controlling the tempering temperature window and time, the carbides are guided to undergo directional self-assembly (alternating with ferrite to form an ordered layered structure), achieving a morphological leap from "granular" to "layered".

[0023] This application follows the principle of "composition-cooling synergistic design." Austenitization ensures complete dissolution of carbides, resulting in austenite with uniform composition. Ultrafast cooling is employed: cooling to room temperature at a rate of ≥50℃ / s. This ultrafast cooling completely avoids the pearlite and bainite transformation regions, directly entering the martensite transformation region to obtain a single lath martensite precursor, precisely constructing the precursor microstructure. The characteristic microstructure of the precursor is as follows: lath martensite width of 200-500 nm; carbon atoms supersaturated dissolved in the α-Fe lattice; high-density dislocations, lath boundaries, and dislocation cell walls; laying the foundation for the self-assembly mechanism of nanolayered structures. The precursor is tempered at 320-380℃ for 3-8 hours, and the microstructure evolution is divided into three stages: Phase 1: Carbon atom segregation and directional nucleation (early tempering stage); Driving force: Chemical driving force (carbon supersaturation) + interfacial energy driving force; Process: Carbon atoms dissolve from the supersaturated α-Fe lattice and segregate towards high-energy regions (lath boundaries, dislocation lines). These high-energy regions become preferential nucleation sites for carbides.

[0024] Key: Due to the presence of a large number of oriented lath boundaries in the precursor, the nucleation sites of carbides exhibit spatial order, laying the foundation for subsequent lamellar growth.

[0025] Phase Two: Preferred growth of carbides and lamellar bonding (mid-tempering). Driving force: crystallographic orientation + elastic strain energy release; Process: Carbides preferentially grow along crystallographic directions that are coherent or semi-coherent with the ferrite matrix (such as {110}α-Fe), forming elongated lamellae. Adjacent carbide lamellae connect with each other to form a continuous lamellae network.

[0026] Key: The tempering temperature (320-380℃) gives carbon atoms a moderate diffusion capacity: enough to support carbide growth without losing directionality due to excessive diffusion. This is the core of achieving a "layered" rather than "granular" structure.

[0027] Stage 3: Laminar coarsening and structural stabilization (late tempering stage); Driving force: The interface can be minimized; Process: Adjacent carbide lamellae further merge and thicken, and the interlamellar spacing tends to be uniform (≤80nm). The ferrite matrix recovers, the dislocation density decreases, and a low-energy interface structure is formed.

[0028] Key point: The tempering time should be controlled within 3-8 hours to ensure that the microstructure transformation is sufficient and orderly, while avoiding excessive coarsening (>8h) or incomplete transformation (<3h).

[0029] Chemical diffusion alone can only produce randomly distributed granular carbides (conventional tempered troostite). Only the synergistic effect of four driving forces (chemical driving force, interfacial energy driving force, elastic strain energy driving force, and crystallographic orientation driving force) can achieve the leap from "random precipitation" to "directional self-assembly" and form a nanolayered structure.

[0030] The chemical driving force originates from the supersaturation of carbon in martensite, providing the thermodynamic driving force for carbide precipitation. The interfacial energy driving force originates from high-density lath boundaries and dislocations, providing the location and orientation for directional nucleation. The elastic strain energy driving force originates from the lattice distortion caused by the martensitic phase transformation, promoting carbide growth along specific directions to release strain. The crystallographic orientation driving force originates from the specific orientation relationship between α-Fe and Fe3C, guiding the preferential growth of carbides to form continuous lamellar layers.

[0031] From a thermodynamic perspective, the morphology of carbides is determined by the principle of minimizing the system's free energy. Granular carbides have lower total interfacial energy but higher elastic strain energy. Lamellar carbides, on the other hand, have higher total interfacial energy but lower elastic strain energy. During conventional tempering, the system tends to minimize interfacial energy, thus leading to carbide spheroidization. This application achieves a dynamic equilibrium between interfacial energy and strain energy by tempering within a specific temperature window (320-380℃), thereby "locking in" the lamellar morphology.

[0032] From a kinetic perspective, the morphology of carbides is controlled by the competition between diffusion rate and interfacial migration rate: below 320℃: the carbon diffusion rate is too low, the carbides are difficult to fully precipitate, and the lamellae are discontinuous; at 320-380℃: the carbon diffusion rate is moderate, the carbide growth rate > the spheroidization rate, and the lamellae are maintained; above 380℃: the carbon diffusion rate is too high, interfacial migration dominates, and the carbides spheroidize rapidly.

[0033] The formation mechanism of the nanolayered structure described in this application has the following unique characteristics: Unique pathway: It completely abandons the traditional "austenite → pearlite" pathway and adopts a brand-new pathway of "lamellar martensite precursor → nanolayered structure".

[0034] The driving force is unique; it does not rely solely on carbon diffusion, but is the result of the synergistic effect of four mechanisms: chemical driving force, interfacial energy driving force, elastic strain energy driving force, and crystallographic orientation driving force.

[0035] Unique kinetic window: Precisely locking onto a specific temperature window of 320-380℃, enabling carbide growth and spheroidization to reach a dynamic balance, achieving a stable existence of the "layered" morphology.

[0036] Uniqueness of structure; obtaining a nanolayered structure with an average interlamellar spacing of ≤80nm and a carbide aspect ratio of ≥5:1, which is different in morphology from traditional pearlite (finer) and conventional tempered structure (layered rather than granular).

[0037] Unique performance characteristics: It achieves an excellent match between tensile strength ≥1300MPa and reduction of area ≥50%, breaking through the traditional bottleneck of high carbon steel having the opposite strength and plasticity. Attached Figure Description

[0038] To more clearly illustrate the technical solutions in the embodiments of this application or the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are only some embodiments of this application. For those skilled in the art, other drawings can be obtained based on the structures shown in these drawings without creative effort.

[0039] Figure 1 This is a microstructure diagram of the high-strength, high-ductility high-carbon steel prepared in Example 1 of this application; Figure 2 The microstructure of the high-strength, high-ductility high-carbon steel prepared in Comparative Example 2 of this application is shown in the diagram. Figure 3 This is a microstructure diagram of the high-strength, high-ductility high-carbon steel prepared in Comparative Example 8 of this application.

[0040] The realization of the purpose, functional features and advantages of this application will be further explained in conjunction with the embodiments and with reference to the accompanying drawings. Detailed implementation mode

[0041] The technical solutions in the embodiments will be described clearly and completely below. Obviously, the described embodiments are only a part of the embodiments of the present application, rather than all the embodiments. All other embodiments obtained by those of ordinary skill in the art based on the embodiments in the present application without making creative efforts belong to the scope of protection of the present application.

[0042] The present application provides a preparation method for high-strength and high-plasticity high-carbon steel, including the following steps: S1. Obtain high-carbon steel, and austenitize the high-carbon steel to obtain austenitic high-carbon steel. By mass percentage, the composition of the high-carbon steel is: C: 0.75% - 0.85%, Mn: 0.5% - 1.0%, Cr: 0.1% - 0.5%, Si: 0.1% - 0.5%, and the rest is Fe and inevitable impurities; The holding temperature for austenitization is 1000 - 1050 °C, and the holding time for austenitization is 20 - 40 min; S2. Rapidly cool the austenitic high-carbon steel to 50 - 150 °C to obtain martensitic high-carbon steel. Among them, the cooling rate of the rapid cooling is ≥50 °C / s, the structure of the martensitic high-carbon steel includes lath martensite, and the content of the lath martensite is ≥90 wt%; The cooling rate is determined according to the equivalent diameter D of the workpiece: when D ≤ 10 mm, the cooling rate is ≥50 °C / s; when 10 mm < D ≤ 20 mm, the cooling rate is ≥80 °C / s; when D > 20 mm, the cooling rate is ≥100 °C / s; the width of the lath martensite is 200 - 500 nm; the structure of the martensitic high-carbon steel includes high-density dislocations, lath boundaries, and dislocation cell walls; S3. Temper the martensitic high-carbon steel to obtain high-strength and high-plasticity high-carbon steel. The tempering temperature is 320 - 380 °C, and the structure of the high-strength and high-plasticity high-carbon steel includes a nano-lamellar structure in which ferrite lamellae and carbide nano-lamellae are alternately arranged; When the tempering temperature is ≥320 °C and ≤350 °C, the tempering time is ≥5 h and ≤8 h; when the tempering temperature is >350 °C and ≤380 °C, the tempering time is ≥3 h and <5 h.

[0043] This application also provides a high-strength, high-ductility, high-carbon steel, comprising: a microstructure comprising alternating ferrite lamellae and carbide nanolamellae, wherein the average interlamellar spacing of the nanolamellae is ≤80nm, and the aspect ratio of the carbide nanolamellae in the nanolamellae is ≥5:1; the high-strength, high-ductility, high-carbon steel has a tensile strength ≥1300MPa, a yield strength ≥1100MPa, a reduction of area ≥50%, and a uniform elongation ≥8.0%.

[0044] The technical solution of this application will be further described below with reference to specific embodiments.

[0045] Example 1 This application provides a method for preparing high-strength, high-ductility high-carbon steel, comprising the following steps: S1. Obtain high carbon steel, and austenitize the high carbon steel to obtain austenitic high carbon steel. The composition of the high carbon steel by mass percentage is: C: 0.80%, Mn: 0.75%, Cr: 0.3%, Si: 0.3%, with the remainder being Fe and unavoidable impurities. The austenitizing holding temperature is 1025℃, and the austenitizing holding time is 30min. S2. Rapidly cooling austenitic high-carbon steel to 100℃ yields martensitic high-carbon steel. The microstructure of the martensitic high-carbon steel includes lath martensite, with a lath martensite content ≥90wt% and a lath martensite width of 200-500nm. The structure of the martensitic high-carbon steel includes high-density dislocations, lath boundaries, and dislocation cell walls. The cooling rate is determined based on the equivalent diameter D of the workpiece: when D=15mm, the cooling rate is ≥80℃ / s. S3. High-strength and high-ductility high-carbon steel is obtained by tempering martensitic high-carbon steel at a temperature of 340℃ for 6 hours.

[0046] The high-strength, high-ductility high-carbon steel prepared in Example 1 was tested, and the results showed that: Please refer to [link / reference needed]. Figure 1 , Figure 1 The microstructure of the high-strength and high-plasticity high-carbon steel prepared in Example 1 of this application is shown in the diagram. The microstructure of the high-strength and high-plasticity high-carbon steel includes a nanolayer structure consisting of alternating ferrite sheets and carbide nanosheets. The average interlamellar spacing of the nanolayer structure is 75 nm, and the aspect ratio of the carbide nanosheets in the nanolayer structure is 8:1. The high-strength and high-carbon steel has a tensile strength of 1350 MPa, a yield strength of 1150 MPa, a reduction of area of ​​52%, and a uniform elongation of 8.5%.

[0047] Example 2 This application provides a method for preparing high-strength, high-ductility high-carbon steel, comprising the following steps: S1. Obtain high carbon steel, and austenitize the high carbon steel to obtain austenitic high carbon steel. The composition of the high carbon steel by mass percentage is: C: 0.85%, Mn: 1.0%, Cr: 0.5%, Si: 0.5%, with the remainder being Fe and unavoidable impurities. The austenitizing holding temperature is 1050℃, and the austenitizing holding time is 20min. S2. Rapidly cooling austenitic high-carbon steel to 150℃ yields martensitic high-carbon steel. The microstructure of the martensitic high-carbon steel includes lath martensite, with a lath martensite content ≥90wt% and a lath martensite width of 200-500nm. The structure of the martensitic high-carbon steel includes high-density dislocations, lath boundaries, and dislocation cell walls. The cooling rate is determined based on the equivalent diameter D of the workpiece: when D=25mm, the cooling rate is ≥100℃. S3. Tempering martensitic high-carbon steel yields high-strength and high-ductility high-carbon steel. The tempering temperature is 380℃ and the tempering time is 3 hours.

[0048] The high-strength and high-carbon steel prepared in Example 2 was tested, and the results showed that the microstructure of the high-strength and high-carbon steel included a nanolayered structure consisting of alternating ferrite lamellae and carbide nanolamellae. The average interlamellae spacing of the nanolayered structure was 72 nm, and the aspect ratio of the carbide nanolamellae in the nanolayered structure was 7:1. The high-strength and high-carbon steel has a tensile strength of 1320 MPa, a yield strength of 1120 MPa, a reduction of area of ​​55%, and a uniform elongation of 9.0%.

[0049] Example 3 This application provides a method for preparing high-strength, high-ductility high-carbon steel, comprising the following steps: S1. Obtain high carbon steel, and austenitize the high carbon steel to obtain austenitic high carbon steel. The composition of the high carbon steel by mass percentage is: C: 0.75%, Mn: 0.5%, Cr: 0.1%, Si: 0.1%, with the remainder being Fe and unavoidable impurities. The austenitizing holding temperature is 1000℃, and the austenitizing holding time is 40min. S2. Rapidly cooling austenitic high-carbon steel to 50-150℃ yields martensitic high-carbon steel. The microstructure of the martensitic high-carbon steel includes lath martensite, with a lath martensite content ≥90wt% and a lath martensite width of 200-500nm. The structure of the martensitic high-carbon steel includes high-density dislocations, lath boundaries, and dislocation cell walls. The cooling rate is determined based on the equivalent diameter D of the workpiece: when D=5mm, the cooling rate is ≥50℃. S3. Tempering martensitic high-carbon steel yields high-strength and high-ductility high-carbon steel. The tempering temperature is 320℃ and the tempering time is 8 hours.

[0050] The high-strength and high-carbon steel prepared in Example 3 was tested, and the results showed that the microstructure of the high-strength and high-carbon steel included a nanolayered structure consisting of alternating ferrite sheets and carbide nanosheets. The average interlamellar spacing of the nanolayered structure was 78 nm, and the aspect ratio of the carbide nanosheets in the nanolayered structure was 6:1. The high-strength and high-carbon steel has a tensile strength of 1330 MPa, a yield strength of 1130 MPa, a reduction of area of ​​51%, and a uniform elongation of 8.0%.

[0051] Comparative Example 1 The difference between this comparative example and Example 1 is that the cooling rate in step S2 is 20°C / s, while the other steps are the same as in Example 1.

[0052] The high-strength and high-plasticity high-carbon steel prepared in Comparative Example 1 was tested. The results showed that the microstructure of the high-strength and high-plasticity high-carbon steel was uneven, with varying lamellar thickness and an average lamellar spacing of 110 nm. The high-strength and high-carbon steel has a tensile strength of 1250 MPa, a yield strength of 1050 MPa, a reduction of area of ​​42%, and a uniform elongation of 6.0%.

[0053] Comparative Example 2 The difference between this comparative example and Example 1 is that the tempering temperature in step S3 is 420°C, while the other steps are the same as in Example 1.

[0054] The high-strength, high-ductility high-carbon steel prepared in Comparative Example 2 was tested, and the results showed that: (Please refer to...) Figure 2 , Figure 2 The microstructure of the high-strength and high-ductility high-carbon steel prepared in Comparative Example 2 of this application is shown. The cementite in the microstructure of the high-strength and high-ductility high-carbon steel exhibits obvious spheroidization and coarsening, and the lamellar structure is severely damaged. The high-strength and high-carbon steel has a tensile strength of 1180 MPa, a yield strength of 1000 MPa, a reduction of area of ​​45%, and a uniform elongation of 6.5%.

[0055] Comparative Example 3 The difference between this comparative example and Example 1 is that the C content of the high-carbon steel in step S1 is 0.60%, while the other steps are the same as in Example 1.

[0056] The high-strength and high-carbon steel prepared in Comparative Example 3 was tested. The results showed that untransformed pearlite and coarse carbides appeared in the microstructure. The nanolayered microstructure was discontinuous and uneven, with an average interlamellar spacing of 120 nm. The high-strength and high-carbon steel has a tensile strength of 1150 MPa, a yield strength of 950 MPa, a reduction of area of ​​38%, and a uniform elongation of 5.5%.

[0057] Comparative Example 4 The difference between this comparative example and Example 1 is that the C content of the high-carbon steel in step S1 is 1.00%, while the other steps are the same as in Example 1.

[0058] The high-strength and high-carbon steel prepared in Comparative Example 4 was tested. The results showed that twinned martensite and coarse undissolved carbides appeared after quenching. After tempering, the carbides were severely coarsened and the lamellar structure was discontinuous. The high-strength and high-carbon steel has a tensile strength of 1280 MPa and a yield strength of 1080 MPa, but a reduction of area of ​​only 32% and a uniform elongation of 4.5%.

[0059] Comparative Example 5 The difference between this comparative example and Example 1 is that the tempering temperature in step S3 is 250°C, while the other steps are the same as in Example 1.

[0060] The high-strength and high-carbon steel prepared in Comparative Example 5 was tested, and the results showed that: the carbon atom diffusion ability was insufficient, the carbide precipitation was incomplete, the nano-layered structure was not formed, and the microstructure was mainly tempered martensite. The tensile strength of high-strength and high-carbon steel is 1420 MPa, but the yield strength is 1180 MPa, the reduction of area is only 28%, and the uniform elongation is 3.5%.

[0061] Comparative Example 6 The difference between this comparative example and Example 1 is that the tempering time in step S3 is 1 hour, while the other steps are the same as in Example 1.

[0062] The high-strength and high-carbon steel prepared in Comparative Example 6 was tested, and the results showed that the carbide precipitation and directional growth were insufficient, the lamellars were discontinuous, and there were a large number of unconverted regions. The high-strength and high-carbon steel has a tensile strength of 1380 MPa, a yield strength of 1160 MPa, a reduction of area of ​​40%, and a uniform elongation of 6.0%.

[0063] Comparative Example 7 The difference between this comparative example and Example 1 is that the tempering time in step S3 is 12 hours, while the other steps are the same as in Example 1.

[0064] The high-strength and high-carbon steel prepared in Comparative Example 7 was tested, and the results showed that the carbides were significantly coarsened, some lamellae fractured and spheroidized, and the average interlamellar spacing was 110 nm. The tensile strength of high-strength and high-carbon steel is reduced to 1220 MPa, the yield strength is 1020 MPa, the reduction of area is 48%, and the uniform elongation is 7.0%.

[0065] Comparative Example 8 High-carbon steel with the same composition as in Example 1 was obtained, heated to 1025°C, held for 30 minutes, and cooled to room temperature using a Steyrmore air-cooling device at an average cooling rate of about 5°C / s to obtain high-strength and high-plasticity high-carbon steel.

[0066] The high-strength, high-ductility high-carbon steel prepared in Comparative Example 8 was tested, and the results show: Please refer to [link / reference needed]. Figure 3 , Figure 3 The image shows the microstructure of the high-strength and high-ductility high-carbon steel prepared in Comparative Example 8 of this application; the microstructure of the high-strength and high-ductility high-carbon steel is sorbite (fine pearlite) with an average lamellar spacing of about 150 nm; The high-strength and high-carbon steel has a tensile strength of 1190 MPa, a yield strength of 980 MPa, a reduction of area of ​​38%, and a uniform elongation of 5.5%.

[0067] Comparative Example 9 High-carbon steel with the same composition as in Example 1 was obtained, heated to 1025°C, held for 30 minutes, oil quenched, and then tempered at 550°C for 2 hours to obtain high-strength and high-plasticity high-carbon steel.

[0068] The high-strength and high-plasticity high-carbon steel prepared in Comparative Example 9 was tested. The results showed that the microstructure of the high-strength and high-plasticity high-carbon steel was tempered sorbite, and the cementite was uniformly distributed in spherical shape with a size of about 100-200 nm. The high-strength and high-carbon steel has a tensile strength of 1250 MPa, a yield strength of 1050 MPa, a reduction of area of ​​48%, and a uniform elongation of 6.0%.

[0069] As can be seen from the above embodiments and comparative examples: Example 1, in conjunction with Comparative Example 1, demonstrates that the cooling rate is crucial for obtaining a uniform lath martensite precursor. When the cooling rate is too low, the transformation of pearlite or bainite cannot be completely suppressed, resulting in an uneven microstructure and a significant decrease in both strength and plasticity. This indicates that a cooling rate ≥50℃ / s is one of the key conditions for obtaining high strength and plasticity; essentially, it ensures that the cooling curve completely avoids the pearlite and bainite transformation regions and directly enters the martensite transformation region.

[0070] Example 1, combined with Comparative Example 2, shows that excessively high tempering temperatures lead to significant spheroidization and coarsening of cementite, disrupting the nanolayered structure and causing the microstructure to degenerate into traditional tempered sorbite (granular), thereby significantly reducing strength and offering limited improvement in plasticity. This confirms that the tempering temperature must be controlled within the 320-380℃ window to avoid the carbide spheroidization rate becoming dominant and to maintain the stability of the lamellar structure.

[0071] Example 1, in conjunction with Comparative Example 3, demonstrates that when the carbon content is too low, the martensite carbon supersaturation is insufficient, the driving force for carbide precipitation during tempering is weakened, and continuous, dense nanosheets cannot be formed. Furthermore, untransformed pearlite remains in the microstructure, resulting in significantly lower strength and plasticity compared to Example 1. This indicates that sufficient carbon content is the material basis for the formation of nanolayered structures.

[0072] Example 1, combined with Comparative Example 4, shows that when the carbon content is too high, twinned martensite and coarse undissolved carbides are formed after quenching. During tempering, the carbides become severely coarsened and unevenly distributed. Although the strength is still acceptable, the plasticity deteriorates sharply. This indicates that the carbon content needs to be controlled within the range of 0.75%-0.85% to balance supersaturation and microstructure uniformity.

[0073] Example 1, combined with Comparative Example 5, shows that when the tempering temperature is too low, the diffusion ability of carbon atoms is insufficient, carbide precipitation is inadequate, the nanolayered structure fails to form, and the microstructure is mainly tempered martensite, exhibiting high strength but extremely low plasticity. This indicates that the temperature needs to be ≥320℃ to activate the directional diffusion and self-assembly process of carbon atoms.

[0074] Example 1, in conjunction with Comparative Example 6, demonstrates that when the tempering time is too short, carbide precipitation and directional growth are insufficient, lamellar layers are discontinuous, and microstructure transformation is incomplete. Although the strength is high, the plasticity is still lower than that of Example 1. This indicates that the tempering time needs to be within the range specified in this application to ensure that the carbides grow sufficiently and connect into continuous lamellar layers.

[0075] Example 1, in conjunction with Comparative Example 7, shows that when the tempering time is too long, the carbides become significantly coarsened, some lamellars fracture and spheroidize, the average interlamellar spacing increases, and the strength decreases significantly. This indicates that the tempering time must be within the range specified in this application to avoid excessive coarsening leading to strength loss.

[0076] Example 1, combined with Comparative Example 8, demonstrates that the nanolayered structure of this application cannot be obtained using the traditional Stellmore air cooling (pearlite phase transformation path). Its structure is sorbite, with significantly lower strength and plasticity than that of Example 1. This proves the significant advantage of the "martensite precursor + intermediate-temperature tempering" path of this application compared to the traditional pearlite path. Its essence lies in a novel mechanism shifting from "high-temperature diffusion-type phase transformation" to "low-temperature shear-type phase transformation + intermediate-temperature precipitation self-assembly."

[0077] Example 1, in conjunction with Comparative Example 9, demonstrates that the traditional "quenching + high-temperature tempering (quenching and tempering)" process yields granular tempered sorbite, which cannot form a nanolayered structure. Although the plasticity is acceptable, the strength is lower than that of Example 1, and the microstructure is completely different from that of this application. This confirms the crucial role of the 320-380℃ tempering window in this application in preserving the lamellar morphology and achieving high strength. Its essence is to achieve a dynamic balance between interfacial energy and strain energy, thereby "locking in" the lamellar morphology rather than spheroidization.

[0078] Examples 1-3, combined with Comparative Examples 1-9, demonstrate that the synergistic process of "composition control + ultra-fast cooling + precise medium-temperature tempering" proposed in this application can successfully prepare a nanolayered structure composed of alternating layers of ferrite and carbide nanosheets. This structure breaks through the performance bottleneck of traditional high-carbon steel, achieving an excellent strength-ductility balance with a tensile strength ≥1300MPa and a reduction of area ≥50%. Its core mechanism lies in the synergistic effect of four driving forces (chemical driving force, interfacial energy driving force, elastic strain energy driving force, and crystallographic orientation driving force) within a specific kinetic window (320-380℃), guiding carbides to undergo "directional self-assembly" from supersaturated martensite, forming a long-range ordered nanolayered structure, thereby achieving a performance combination that cannot be achieved by traditional pearlite processing and tempering processes.

[0079] The above description is only a preferred embodiment of this application and does not limit the patent scope of this application. All equivalent structural transformations made using the content of this application's specification under the inventive concept of this application, or direct / indirect applications in other related technical fields, are included within the patent protection scope of this application.

Claims

1. A method for producing a high-strength plastic high-carbon steel, characterized by, It includes the following steps: S1. Obtain high-carbon steel, and subject the high-carbon steel to austenitization to obtain austenitic high-carbon steel. By mass percentage, the composition of the high-carbon steel is: C: 0.75% - 0.85%, Mn: 0.5% - 1.0%, Cr: 0.1% - 0.5%, Si: 0.1% - 0.5%, and the balance is Fe and unavoidable impurities; S2. Rapidly cool the austenitic high-carbon steel to 50 - 150°C to obtain martensitic high-carbon steel. Among them, the cooling rate of the rapid cooling is ≥50°C / s, the structure of the martensitic high-carbon steel includes lath martensite, and the content of the lath martensite is ≥90wt%; S3. Temper the martensitic high-carbon steel to obtain high-strength and high-ductility high-carbon steel. The temperature of the tempering is 320 - 380°C, and the structure of the high-strength and high-ductility high-carbon steel includes a nano-laminated structure in which ferrite lamellae and carbide nano-lamellae are alternately arranged.

2. The method of claim 1, wherein the high-strength ductile high-carbon steel is prepared by the steps of: In step S1, the holding temperature of the austenitization is 1000 - 1050°C, and the holding time of the austenitization is 20 - 40 min. ​ 3. The method of claim 1, wherein the high-strength ductile high-carbon steel is prepared by the steps of: In step S2, the cooling rate is determined according to the equivalent diameter D of the workpiece: when D ≤ 10 mm, the cooling rate is ≥50°C / s; when 10 mm < D ≤ 20 mm, the cooling rate is ≥80°C / s; when D > 20 mm, the cooling rate is ≥100°C / s. ​ 4. The method for preparing high-strength, high-ductility high-carbon steel according to claim 1, characterized in that, In step S3, when the temperature of the tempering is ≥320°C and ≤350°C, the tempering time is ≥5 h and ≤8 h.

5. The method for preparing high-strength, high-ductility, high-carbon steel according to claim 1, characterized in that, In step S3, when the temperature of the tempering is > 350°C and ≤380°C, the tempering time is ≥3 h and < 5 h.

6. The method for preparing high-strength, high-ductility high-carbon steel according to claim 1, characterized in that, In step S2, the width of the lath martensite is 200 - 500 nm.

7. The method for preparing high-strength, high-ductility, high-carbon steel according to claim 1, characterized in that, In step S2, the structure of the martensitic high-carbon steel includes high-density dislocations, lath boundaries, and dislocation cell walls.

8. A high-strength, high-ductility, high-carbon steel, characterized in that, It is prepared by using the preparation method of a high-strength and high-ductility high-carbon steel according to any one of claims 1 - 7.

9. The high-strength, high-ductility, high-carbon steel according to claim 8, characterized in that, The structure of the high-strength and high-ductility high-carbon steel includes a nano-laminated structure in which ferrite lamellae and carbide nano-lamellae are alternately arranged. The average lamellar spacing of the nano-laminated structure is ≤80 nm, and the aspect ratio of the carbide nano-lamellae in the nano-laminated structure is ≥5:

1.

10. A high-strength, high-ductility, high-carbon steel according to claim 8, characterized in that, The tensile strength of the high-strength and high-ductility high-carbon steel is ≥1300 MPa, the yield strength of the high-strength and high-ductility high-carbon steel is ≥1100 MPa, the reduction of area of the high-strength and high-ductility high-carbon steel is ≥50%, and the uniform elongation of the high-strength and high-ductility high-carbon steel is ≥8.0%.

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