High strength and plasticity cord steel and its preparation method

Abstract

The application belongs to the technical field of cord steel preparation, and particularly relates to a high-strength and high-plasticity cord steel and a preparation method thereof, the preparation method comprising the following steps: S1, obtaining a cord steel casting blank, heating, rolling and cooling the cord steel casting blank to obtain a cord steel wire rod, and the structure of the cord steel wire rod is a pearlite structure with an average interlamellar spacing of 95-100 nm; S2, cold drawing the cord steel wire rod to obtain a cord steel wire material, and the total surface reduction of the cold drawing is greater than or equal to 85%; and S3, heat treating the cord steel wire material to obtain the high-strength and high-plasticity cord steel, and the holding temperature of the heat treatment is 280-300 DEG C. According to the application, a pearlite matrix is obtained through alloy composition and cooling design, and severe plastic deformation is performed to construct an initial state with high defect density, low-temperature heat treatment is adopted to drive short-range diffusion of interface carbon atoms to construct a composition gradient and drive diffusion of micro-alloy elements to regulate nanometer precipitates, and the interface plasticization and precipitation reinforcement are synergized.

Description

Technical Field

[0001] This application relates to the field of cord steel preparation technology, specifically, this application relates to a high-strength and ductile cord steel and its preparation method. Background Technology

[0002] Ultra-high strength pearlitic cord steel (typical grades such as SWRS82B and 92B) is a key basic material for manufacturing high-grade tire cords, special steel wire ropes, and prestressed steel strands. Its core performance requirements are to achieve the highest possible tensile strength (typically ≥2000MPa) and sufficient plasticity reserve (elongation after fracture) while ensuring excellent torsional and bending fatigue performance. Through the classic process of "hot-rolled wire rod → deep cold drawing," the material can easily achieve the strength target through severe work hardening. However, with the accumulation of cold deformation (reduction of area), the plasticity of pearlitic steel decreases sharply, and the elongation after fracture is usually less than 3%, exhibiting a typical "high strength - low plasticity" imbalance. This brittle tendency makes the material exceptionally sensitive to notches and stress concentrations during subsequent stranding, cabling, and dynamic service, easily leading to early fracture and severely limiting its application in high-reliability scenarios.

[0003] To address this challenge of the inversion of strong plasticity, the industry has explored solutions from two main perspectives: Existing technologies focus on microstructure control and alloying. For example, traditional production often relies on the Stelmor air-cooling process, controlling cooling (typically at a rate of 3-10°C / s) to obtain a pearlitic microstructure. However, pearlite formed at conventional cooling rates has a relatively coarse average interlamellar spacing (typically >150nm) and may be accompanied by the precipitation of proeutectoid ferrite. Attempts to refine the microstructure by increasing the cooling rate have been made, but without precise phase transformation kinetics guidance, excessively rapid cooling can cause the material to enter the bainite region of the continuous cooling transformation (CCT) curve, forming a hard and brittle mixed bainite and pearlite microstructure, which negatively impacts overall performance. Furthermore, simple alloying (such as simply increasing the Mn and Si content) can strengthen ferrite or refine grains to some extent, but its effect on refining the interlamellar spacing is limited and may worsen the drawing performance of the wire rod and increase costs.

[0004] Existing technologies rely on post-processing modification, with heat treatment after cold drawing being the most common method to improve plasticity. For example, traditional processes often employ medium-to-high temperature (350-500℃) annealing. The mechanism involves softening the ferrite matrix through recovery and recrystallization, and promoting cementite lamellarization, thereby releasing internal stress and improving ductility. However, this significant change in macrostructure inevitably leads to a substantial drop in strength (strength loss can reach 15%-30%), negating its fundamental value as an ultra-high strength material. Another approach is low-temperature tempering (200-300℃), aiming to eliminate some internal stress without significantly altering the microstructure. However, for cord steel prepared using traditional compositions and processes, the improvement in plasticity from simple low-temperature tempering is very limited (usually only 1%-2%), failing to meet the requirements for high plasticity.

[0005] Delving into the essence of its microscopic deformation, the root cause of the plasticity bottleneck in ultra-high strength pearlitic cord steel lies in its nanoscale two-phase (ferrite and cementite) interface. After deep cold drawing, a sharp, high-energy, and fragile interface forms between the two phases. Under external force, dislocations accumulate severely at this interface, leading to extremely high local stress concentration, making microcracks easily initiate and propagate along the interface. Therefore, the core of improving plasticity lies in reconstructing this interface, transforming it from a "brittle barrier" into a "flexible buffer layer" capable of coordinating deformation and transmitting dislocations. However, existing technologies face a fundamental dilemma when optimizing this key interface. If sufficiently high energy (such as medium to high temperatures) is used to drive the full diffusion and reconstruction of interface atoms, it will inevitably trigger the recrystallization of the matrix ferrite and the spheroidization of cementite simultaneously, leading to the collapse of the strengthening structure; if the processing temperature is too low, the diffusion dynamics of interface atoms are insufficient, making effective interface modification impossible. The underlying reason lies in the failure to treat "atomic-level precise control of interface structure" and "overall maintenance of matrix strengthening structure" as two independent dimensions that can be designed separately and then synergized. Current technology lacks a cascade solution that can first prefabricate a fine matrix structure with both "interface tunability" and "strengthened stability," and then selectively modify only the interface region through a precise, directional post-processing step. More critically, the catalytic effect of high-density crystal defects (dislocations, interfaces) introduced by cold deformation as rapid channels for atomic diffusion on subsequent low-temperature interface engineering has not been fully recognized and utilized. Furthermore, the potential synergistic relationship between the carbonitride precipitation behavior of microalloying elements (such as V and Cr) and the phase interface diffusion process has not been systematically integrated and utilized in the process design.

[0006] Therefore, the current technological gap lies in the need for a novel, systematic material design and preparation method based on phase transformation and precipitation kinetics. This method should be able to: obtain a pure, ultrafine pearlite matrix with microalloying elements in a metastable supersaturated state through precise alloy design and cooling control (based on CCT / TTT curves); apply severe plastic deformation to this matrix to construct an initial state with high defect density; design a low-temperature heat treatment window that utilizes the diffusion channels provided by the pre-deformation to simultaneously and selectively drive short-range diffusion of interfacial carbon atoms to construct a compositional gradient and the diffusion of microalloying elements to regulate the nano-precipitated phase. Ultimately, without affecting the macroscopic morphology of the matrix, achieve a synergistic effect of interfacial plasticity and precipitation reinforcement, fundamentally breaking the traditional trade-off between strength and plasticity. This application aims to fill this technological gap. Summary of the Invention

[0007] To solve the above-mentioned technical problems, this application provides a method for preparing high-strength and high-ductility cord steel, comprising the following steps: S1, obtaining a cord steel billet, heating, rolling, and cooling the cord steel billet to obtain cord steel wire, wherein the microstructure of the cord steel wire is a pearlitic microstructure with an average interlamellar spacing of 95-100 nm; the composition of the cord steel billet by mass percentage is: C: 0.80%-0.85%, Si: 0.15%-0.30%, Mn: 0.40%-0.85%. S1. 70%, Cr: 0.15%-0.30%, V: 0.05%-0.15%, N: 0.0080%-0.0120%, P≤0.012%, S≤0.008%, the remainder being Fe and unavoidable impurities; S2. The cord steel wire is cold-drawn to obtain cord steel wire, the total area reduction of the cold drawing being ≥85%; S3. The cord steel wire is heat-treated to obtain high-strength plastic cord steel, the heat treatment holding temperature being 280-300℃.

[0008] In a preferred embodiment of the method for preparing high-strength and ductile cord steel described in this application, the heating temperature in step S1 is 1020-1080℃.

[0009] As a preferred embodiment of the preparation method of high-strength and ductile cord steel described in this application, in step S1, the cooling method is specifically as follows: cooling to 480-520℃ at a cooling rate of 8-12℃ / s and holding at that temperature for 10-30 seconds.

[0010] In a preferred embodiment of the method for preparing high-strength and ductile cord steel described in this application, the cold drawing process in step S2 consists of 3-5 passes.

[0011] In a preferred embodiment of the method for preparing high-strength and ductile cord steel described in this application, the heat treatment holding time in step S3 is 30-90 min.

[0012] As a preferred embodiment of the preparation method of high-strength and ductile cord steel described in this application, in step S3, when the temperature of the heat treatment is ≥280℃ and ≤290℃, the heat treatment holding time is ≥60min and ≤90min.

[0013] As a preferred embodiment of the preparation method of high-strength and ductile cord steel described in this application, in step S3, when the temperature of the heat treatment is >290℃ and ≤300℃, the heat treatment holding time is ≥30min and <60min.

[0014] This application also provides a high-strength, high-ductility cord steel, which is prepared using the above-described method for preparing high-strength, high-ductility cord steel.

[0015] As a preferred embodiment of the high-strength plastic cord steel described in this application, the high-strength plastic cord steel has a tensile strength ≥2000MPa, a yield strength ≥1800MPa, an elongation ≥8.5%, and a reduction of area ≥40%.

[0016] As a preferred embodiment of the high-strength and ductile cord steel described in this application, the microstructure of the high-strength and ductile cord steel includes: pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite is ≤95nm. The pearlite is composed of cementite and ferrite, wherein carbon atoms diffuse from the cementite to the ferrite, forming a carbon concentration gradient transition region with a width of 50-100nm at the interface between the cementite and the ferrite. The average size of the (Cr,V)CN composite precipitates is ≤10nm, and the (Cr,V)CN composite precipitates are dispersedly distributed in the ferrite.

[0017] The beneficial effects of this application are as follows: This application proposes a method for preparing high-strength and ductile cord steel. This application obtains a pearlite matrix through the synergistic design of alloy composition and cooling, and constructs an initial state with high defect density through severe plastic deformation. Low-temperature heat treatment is used to utilize the diffusion channels provided by the pre-deformation to drive the short-range diffusion of carbon atoms at the interface to construct a composition gradient, and to drive the diffusion of microalloying elements to regulate the nano-precipitated phase, thereby achieving the synergistic effect of interface plasticization and precipitation reinforcement.

[0018] This application utilizes optimal undercooling to obtain a pure pearlitic microstructure, rapidly traversing the high-temperature precipitation range of Cr and V carbonitrides in the austenite region, maximally retaining them within the supersaturated solid solution, thus creating conditions for subsequent nanoprecipitation. According to CCT curves, a cooling rate of 8-12 °C / s corresponds to the fastest pearlitic transformation kinetics and effectively avoids the "safe window" of the proeutectoid ferrite and bainite transformation regions. The undercooling (ΔT) provided by this cooling rate can stably control the interlamellar spacing at 95-100 nm using the Zener-Hillert model (λ∝1 / ΔT). Short-duration isothermal treatment at 480-520 °C ensures complete phase transformation and avoids residual austenite. This cooling regime suppresses the nucleation and growth of coarse (Cr, V) CN at grain boundaries at temperatures above 600 °C. At lower phase transformation temperatures, the diffusion capacity of Cr and V atoms is limited, thus promoting the formation of a large number of dispersed (Cr,V)CN composite nanoprecipitates with a size ≤10 nm in ferrite during or after the phase transformation. Cr increases the eutectoid point (e.g., Ac1) temperature, expands the process window for pearlite transformation, and at the same time reduces the diffusion coefficient of carbon in austenite. Combined with the accelerated cooling effect, this has a stronger inhibitory effect on lamellar growth.

[0019] This application, based on the high dislocation density microstructure introduced by cold deformation, utilizes selective diffusion at low temperatures to simultaneously and independently optimize the phase interface structure and control the precipitate state, achieving a significant improvement in plasticity and precise maintenance or slight increase in strength. The high-density dislocations and phase interfaces generated by cold drawing provide carbon atoms with rapid short-circuit diffusion channels. At 280-300℃, carbon atoms possess sufficient activity for short-range diffusion, diffusing from cementite lamellars to adjacent ferrite, forming a "flexible interface layer" with a continuous compositional transition. This gradient region effectively coordinates the strain incompatibility during the deformation of the two phases, significantly reducing the resistance of dislocations crossing the interface, thereby increasing the elongation from approximately 2% in the cold-drawn state to ≥8.5%. This temperature range is precisely the optimal temperature for V and Cr atoms to undergo limited diffusion, driving the existing nano-(Cr,V)CN composite precipitates to evolve into a more stable state through Ostwald ripening. This process makes the precipitate distribution more uniform and stable, producing a significant coherent or semi-coherent precipitation strengthening effect, sufficient to precisely compensate for the slight strength loss that may be caused by interface relaxation. Undeformed pearlite, under the same low-temperature treatment, exhibits extremely low diffusion efficiency between carbon and alloy atoms, making it impossible to form an effective gradient and regulate precipitation.

[0020] In the technical solution described in this application, the average interlamellar spacing of pearlite exhibits a regular variation across different process stages. During the cold drawing stage, the large plastic deformation with a total area reduction of ≥85% causes the pearlite microstructure to undergo severe deformation under the combined action of axial tensile stress and radial compressive stress. The refinement of the interlamellar spacing stems from the synergistic effect of geometry and crystallography. Initially randomly oriented pearlite clusters gradually rotate, causing their lamellar normals to tend towards perpendicularity to the drawing axis, thus significantly reducing the apparent interlamellar spacing measured on the longitudinal section. Simultaneously, the hard and brittle cementite lamellars bend, twist, or even fracture under shear stress, resulting in a shortened length of the fractured cementite fragments and a corresponding decrease in the effective distance between adjacent cementite lamellars. Furthermore, the ferrite lamellars, possessing sufficient slip systems, undergo a large amount of dislocation slip, being thinned along the thickness direction. The combined effect of these three factors further refines the pearlite with an initial interlamellar spacing of 95-100 nm to 85-95 nm, a refinement of approximately 5-15%. In the subsequent low-temperature heat treatment stage of 280-300℃, the temperature is far below the 400℃ threshold required for significant spheroidization of cementite. The long-range diffusion ability of iron and carbon atoms is extremely low, which is insufficient to drive the dissolution, fracture or spheroidization of cementite plates. This temperature range only causes partial recovery of the high-density dislocations introduced by cold deformation, eliminating some internal stress. At the same time, although carbon atoms have the ability to migrate from the cementite surface to the adjacent ferrite to form a concentration gradient, it will not cause the overall thinning of cementite plates. The limited Ostwald ripening of the (Cr,V)CN composite precipitate phase only occurs inside the ferrite and is unrelated to the cementite plates. Therefore, low-temperature heat treatment will not cause plate coarsening or spheroidization.

[0021] The high-strength, high-ductility cord steel prepared in this application has the following characteristics: The high-strength and high-ductility cord steel has a tensile strength ≥2000MPa, a yield strength ≥1800MPa, an elongation ≥8.5%, and a reduction of area ≥40%. The microstructure of the high-strength and high-ductility cord steel includes pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite is ≤95nm. The pearlite is composed of cementite and ferrite, wherein carbon atoms diffuse from the cementite to the ferrite, forming a carbon concentration gradient transition region with a width of 50-100nm at the interface between the cementite and the ferrite. The average size of the (Cr,V)CN composite precipitates is ≤10nm, and the (Cr,V)CN composite precipitates are dispersed within the ferrite. Attached Figure Description

[0022] 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.

[0023] Figure 1 This is a diagram of the pearlitic lamellar structure of the high-strength and ductile cord steel prepared in Example 1 of this application; Figure 2 This is a diagram of the precipitated phase structure of the high-strength, high-ductility cord steel prepared in Example 1 of this application.

[0024] 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

[0025] The technical solutions in the embodiments will be clearly and completely described below. Obviously, the described embodiments are only some embodiments of this application, and not all embodiments. Based on the embodiments in this application, all other embodiments obtained by those of ordinary skill in the art without creative effort are within the scope of protection of this application.

[0026] This application provides a method for preparing high-strength, high-ductility steel cord, comprising the following steps: S1. Obtain a cord steel billet, heat, roll, and cool the cord steel billet to obtain cord steel wire. The microstructure of the cord steel wire is pearlite with an average interlamellar spacing of 95-100 nm. By mass percentage, the composition of the cord steel billet is: C: 0.80%-0.85%, Si: 0.15%-0.30%, Mn: 0.40%-0.70%, Cr: 0.15%-0.30%, V: 0.05%-0.15%, N: 0.0080%-0.0120%, P≤0.012%, S≤0.008%, with the remainder being Fe and unavoidable impurities. The heating temperature is 1020-1080℃; the cooling method is specifically: cooling to 480-520℃ at a cooling rate of 8-12℃ / s and holding at that temperature for 10-30 seconds; S2. The cord steel wire is cold-drawn to obtain cord steel wire, wherein the total area shrinkage of the cold drawing is ≥85%; the number of cold drawing passes is 3-5. S3. The cord steel wire is heat-treated to obtain high-strength ductile steel cord. The heat treatment holding temperature is 280-300℃, and the heat treatment holding time is 30-90min. When the temperature of the heat treatment is ≥280℃ and ≤290℃, the holding time of the heat treatment is ≥60min and ≤90min; when the temperature of the heat treatment is >290℃ and ≤300℃, the holding time of the heat treatment is ≥30min and <60min.

[0027] This application also provides a high-strength and high-ductility cord steel, comprising: a tensile strength ≥2000MPa, a yield strength ≥1800MPa, an elongation ≥8.5%, and a reduction of area ≥40%; the microstructure of the high-strength and high-ductility cord steel comprises: pearlite and (Cr,V)CN composite precipitates, wherein the average interlamellar spacing of the pearlite is ≤95nm, the pearlite is composed of cementite and ferrite, wherein carbon atoms diffuse from the cementite to the ferrite, forming a carbon concentration gradient transition region with a width of 50-100nm at the interface between the cementite and the ferrite, the average size of the (Cr,V)CN composite precipitates is ≤10nm, and the (Cr,V)CN composite precipitates are dispersedly distributed in the ferrite.

[0028] The technical solution of this application will be further described below with reference to specific embodiments.

[0029] In this application, the average interlamellar spacing of pearlite is determined using the following method: Samples are taken from the cross-section of cord steel wire, cord steel filament, or high-strength ductile cord steel. After mechanical polishing and electrolytic double-spray etching with a 10% nitric acid alcohol solution, the microstructure of the sample is observed using a transmission electron microscope. At least five different pearlite clusters are randomly selected from the core, half radius, and edge of each sample. The interlamellar spacing is measured using the cross-section method (drawing a straight line perpendicular to the lamellar direction, counting the number of cementite lamellars passed through by the line, and dividing the line length by the number of lamellars). A total of no less than 50 interlamellar spacing values ​​are measured for each sample, and the arithmetic mean is taken as the average interlamellar spacing of the sample.

[0030] Example 1 This application provides a method for preparing high-strength, high-ductility steel cord, comprising the following steps: S1. Obtain the cord steel billet, heat, roll, and cool the cord steel billet to obtain cord steel wire. The microstructure of the cord steel wire is pearlite with an average interlamellar spacing of 97 nm. By mass percentage, the composition of the cord steel billet is: C: 0.82%, Si: 0.20%, Mn: 0.55%, Cr: 0.25%, V: 0.10%, N: 0.0100%, P≤0.012%, S≤0.008%, with the remainder being Fe and unavoidable impurities. The heating temperature is 1050℃; the cooling method is as follows: cool to 500℃ at a cooling rate of 10℃ / s and hold for 20 seconds; S2. The cord steel wire is cold-drawn to obtain cord steel wire. The total area reduction of the cold drawing is 86%. The number of cold drawing passes is 4. After cold drawing, the average interlamellar spacing of pearlite is refined to 90nm. S3. High-strength and ductile cord steel is obtained by heat treatment of the cord steel wire. The heat treatment holding temperature is 290℃ and the heat treatment holding time is 60min. After low-temperature heat treatment, the average interlamellar spacing of pearlite is maintained at 90nm. The high-strength, high-ductility cord steel prepared in Example 1 was tested. The results showed that the high-strength, high-ductility cord steel had a tensile strength of 2090 MPa, a yield strength of 1862 MPa, an elongation of 9.2%, and a reduction of area of ​​48%. The microstructure of the high-strength, high-ductility cord steel included pearlite and (Cr, V)CN composite precipitates. Please refer to [link to relevant documentation]. Figure 1 , Figure 1 This is a diagram of the pearlitic lamellar structure of the high-strength, high-ductility cord steel prepared in Example 1 of this application. The average interlamellar spacing of the pearlite is 90 nm. The pearlite consists of cementite and ferrite, wherein carbon atoms diffuse from cementite to ferrite, forming a carbon concentration gradient transition region with a width of 58 nm at the interface between cementite and ferrite. Please refer to [link to relevant documentation]. Figure 2 , Figure 2 The diagram shows the precipitate structure of the high-strength and ductile cord steel prepared in Example 1 of this application; the average size of the (Cr,V)CN composite precipitate is 7.1 nm, and the (Cr,V)CN composite precipitate is dispersed in the ferrite.

[0031] Example 2 This application provides a method for preparing high-strength, high-ductility steel cord, comprising the following steps: S1. Obtain the cord steel billet, heat, roll, and cool the cord steel billet to obtain cord steel wire. The microstructure of the cord steel wire is pearlite with an average interlamellar spacing of 99 nm. By mass percentage, the composition of the cord steel billet is: C: 0.85%, Si: 0.30%, Mn: 0.70%, Cr: 0.30%, V: 0.15%, N: 0.0120%, P≤0.012%, S≤0.008%, with the remainder being Fe and unavoidable impurities. The heating temperature is 1080℃; the cooling method is as follows: cool to 520℃ at a cooling rate of 12℃ / s and hold for 10 seconds; S2. The cord steel wire is cold-drawn to obtain cord steel wire. The total area reduction of the cold drawing is 88%. The number of cold drawing passes is 5. After cold drawing, the average interlamellar spacing of pearlite is refined to 92nm. S3. High-strength and ductile cord steel is obtained by heat treatment of the cord steel wire. The heat treatment temperature is 300℃ and the heat treatment time is 30min. After low-temperature heat treatment, the average interlamellar spacing of pearlite is maintained at 92nm. The high-strength and ductile cord steel prepared in Example 2 was tested. The results showed that the high-strength and ductile cord steel had a tensile strength of 2120 MPa, a yield strength of 1800 MPa, an elongation of 8.6%, and a reduction of area of ​​45%. The microstructure of the high-strength and ductile cord steel included pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite was 92 nm. The pearlite was composed of cementite and ferrite. Carbon atoms diffused from cementite to ferrite, forming a carbon concentration gradient transition region with a width of 80 nm at the interface between cementite and ferrite. The average size of the (Cr,V)CN composite precipitates was 6.9 nm. The (Cr,V)CN composite precipitates were dispersed in the ferrite.

[0032] Example 3 This application provides a method for preparing high-strength, high-ductility steel cord, comprising the following steps: S1. Obtain the cord steel billet, heat, roll, and cool the cord steel billet to obtain cord steel wire. The microstructure of the cord steel wire is pearlite with an average interlamellar spacing of 95 nm. By mass percentage, the composition of the cord steel billet is: C: 0.80%, Si: 0.15%, Mn: 0.40%, Cr: 0.15%, V: 0.05%, N: 0.0080%, P≤0.012%, S≤0.008%, with the remainder being Fe and unavoidable impurities. The heating temperature is 1020℃; the cooling method is as follows: cooling to 480℃ at a cooling rate of 8℃ / s and holding at that temperature for 30 seconds; S2. The cord steel wire is cold-drawn to obtain cord steel wire. The total area reduction of the cold drawing is 85%. The cold drawing is performed in 3 passes. After cold drawing, the average interlamellar spacing of pearlite is refined to 88nm. S3. High-strength and ductile cord steel is obtained by heat treatment of the cord steel wire. The heat treatment temperature is 280℃ and the heat treatment time is 90min. After low-temperature heat treatment, the average interlamellar spacing of pearlite is maintained at 88nm. The high-strength and ductile cord steel prepared in Example 3 was tested. The results showed that the high-strength and ductile cord steel had a tensile strength of 2040 MPa, a yield strength of 1816 MPa, an elongation of 8.7%, and a reduction of area of ​​43%. The microstructure of the high-strength and ductile cord steel included pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite was 88 nm. The pearlite was composed of cementite and ferrite. Carbon atoms diffused from the cementite to the ferrite, forming a carbon concentration gradient transition region with a width of 65 nm at the interface between the cementite and ferrite. The average size of the (Cr,V)CN composite precipitates was 10 nm. The (Cr,V)CN composite precipitates were dispersed in the ferrite.

[0033] Comparative Example 1 The difference between this comparative example and Example 1 is that step S3 is omitted, while the other steps are the same as in Example 1.

[0034] The cord steel wire prepared in Comparative Example 1 was tested. The results showed that the tensile strength of the cord steel wire was 2180 MPa, the yield strength was 1990 MPa, the elongation was 2.3%, and the reduction of area was 18%. The microstructure of the cord steel wire included pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite was 90 nm. The pearlite was composed of cementite and ferrite. The interface between cementite and ferrite was sharp, and there was no obvious carbon concentration gradient transition zone. The average size of the (Cr,V)CN composite precipitates was 7.5 nm, but the distribution was uneven.

[0035] Comparative Example 2 The difference between this comparative example and Example 1 is that the composition of the cord steel billet is: C: 0.82%, Si: 0.20%, Mn: 0.55%, P≤0.012%, S≤0.008%, with the remainder being Fe and unavoidable impurities. The other steps are the same as in Example 1.

[0036] The high-strength and ductile cord steel prepared in Comparative Example 2 was tested. The results showed that the high-strength and ductile cord steel had a tensile strength of 1895 MPa, a yield strength of 1720 MPa, an elongation of 6.1%, and a reduction of area of ​​32%. The microstructure of the cord steel wire included pearlite with an average interlamellar spacing of 115 nm. The pearlite was composed of cementite and ferrite, with no (Cr, V)CN composite precipitates. The carbon concentration gradient at the cementite-ferrite interface was not obvious and its width was less than 10 nm.

[0037] Comparative Example 3 The difference between this comparative example and Example 1 is that the cooling rate in step S1 is 3°C / s, while the other steps are the same as in Example 1.

[0038] The high-strength and ductile cord steel prepared in Comparative Example 3 was tested. The results showed that the high-strength and ductile cord steel had a tensile strength of 2010 MPa, a yield strength of 1780 MPa, an elongation of 4.5%, and a reduction of area of ​​28%. The microstructure of the cord steel wire included pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite was 170 nm, and a small amount of proeutectoid ferrite was present. The average size of the (Cr,V)CN composite precipitates was 25 nm, and the distribution was uneven.

[0039] Comparative Example 4 The difference between this comparative example and Example 1 is that the heat treatment holding temperature in step S3 is 400°C, while the other steps are the same as in Example 1.

[0040] The high-strength and ductile cord steel prepared in Comparative Example 4 was tested. The results showed that the high-strength and ductile cord steel had a tensile strength of 1850 MPa, a yield strength of 1620 MPa, an elongation of 6.8%, and a reduction of area of ​​38%. The microstructure of the cord steel wire included pearlite and (Cr,V)CN composite precipitates. However, the cementite underwent significant spheroidization, the ferrite underwent recrystallization, the interfacial carbon concentration gradient disappeared, the lamellar structure was destroyed, and the effective interlamellar spacing could not be measured. The average size of the (Cr,V)CN composite precipitates was 20 nm.

[0041] Comparative Example 5 The difference between this comparative example and Example 1 is that the heat treatment in step S3 is performed first, followed by the cold drawing in step S2. The other steps are the same as in Example 1.

[0042] The cord steel wire prepared in Comparative Example 5 was tested. The results showed that the tensile strength of the cord steel wire was 1920 MPa, the yield strength was 1700 MPa, the elongation was 3.8%, and the reduction of area was 22%. The microstructure of the cord steel wire included pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite was 97 nm. However, due to the lack of introduction of high-density dislocation channels after cold drawing, no obvious carbon concentration gradient transition zone was formed at the interface between cementite and ferrite. The average size of the (Cr,V)CN composite precipitates was 12 nm.

[0043] Comparative Example 6 The difference between this comparative example and Example 1 is that the total area reduction of cold drawing in step S2 is 60%, while the other steps are the same as in Example 1.

[0044] The high-strength and ductile cord steel prepared in Comparative Example 6 was tested. The results showed that the high-strength and ductile cord steel had a tensile strength of 1980 MPa, a yield strength of 1750 MPa, an elongation of 4.2%, and a reduction of area of ​​25%. The microstructure of the cord steel wire included pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite was 95 nm. The pearlite was composed of cementite and ferrite. Carbon atoms diffused from cementite to ferrite, forming a carbon concentration gradient transition region with a width of 15 nm at the interface between cementite and ferrite. The average size of the (Cr,V)CN composite precipitates was 8 nm. The (Cr,V)CN composite precipitates were dispersed in the ferrite.

[0045] As can be seen from the above embodiments and comparative examples: Example 1, in conjunction with Comparative Example 1, shows that although Comparative Example 1, which did not undergo the low-temperature heat treatment in step S3, had slightly higher tensile strength and yield strength than Example 1, its elongation and reduction of area were significantly lower. This indicates that without low-temperature heat treatment, the cold-drawn cord steel wire, while possessing high strength, exhibits extremely poor plasticity, failing to meet the plasticity requirements for cord steel applications. The low-temperature heat treatment of 280-300℃ in this application can significantly improve elongation and reduction of area while maintaining high strength, achieving a good balance between strength and plasticity.

[0046] Example 1, combined with Comparative Example 2, shows that Comparative Example 2, without the addition of Cr, V, and N elements, exhibits significantly lower tensile strength, yield strength, elongation, and reduction of area compared to Example 1. Furthermore, it lacks (Cr, V)CN composite precipitates, and the average interlamellar spacing of the pearlite is slightly coarser. This indicates that the synergistic addition of Cr, V, and N can further refine the average interlamellar spacing of the pearlite and form a dispersed nanoscale (Cr, V)CN composite precipitate in the ferrite, resulting in a significant precipitation strengthening effect. This provides a favorable microstructure basis for the construction of interface gradients in subsequent low-temperature heat treatment.

[0047] Example 1, in conjunction with Comparative Example 3, demonstrates that reducing the cooling rate to 3°C / s in Comparative Example 3 resulted in an increase in the average interlamellar spacing of pearlite and the appearance of proeutectoid ferrite. The coarsening of the (Cr,V)CN composite precipitate led to significantly lower tensile strength, elongation, and reduction of area compared to Example 1. This indicates that a rapid cooling rate of 8-12°C / s is a key process condition for obtaining a pearlite microstructure with an average interlamellar spacing ≤100 nm, suppressing proeutectoid ferrite precipitation, and maintaining a fine and dispersed distribution of the (Cr,V)CN composite precipitate.

[0048] Example 1, in conjunction with Comparative Example 4, shows that while increasing the heat treatment temperature to 400℃ in Comparative Example 4 improved elongation and reduction of area compared to Comparative Example 1, it significantly reduced tensile strength and yield strength, resulting in a high strength loss rate. Microstructural observation revealed marked spheroidization of cementite, recrystallization of ferrite, and the disappearance of the interfacial carbon concentration gradient. This indicates that 280-300℃ is the critical temperature window for achieving "selective interfacial modification without triggering matrix degradation" in the technical solution of this application; temperatures exceeding this range will lead to the collapse of the strengthened structure.

[0049] Example 1, in conjunction with Comparative Example 5, demonstrates that Comparative Example 5, by reversing the order of heat treatment and cold drawing, exhibits significantly lower tensile strength, yield strength, elongation, and reduction of area compared to Example 1. Microstructural analysis reveals that, due to the failure to form high-density dislocation channels after cold drawing, even subsequent low-temperature heat treatment prevents effective diffusion of carbon atoms into the ferrite to form a sufficiently wide carbon concentration gradient transition region. This indicates that the process sequence of this application—"introducing high defect density through cold drawing followed by low-temperature heat treatment to drive interfacial diffusion using defect channels"—possesses irreplaceable technical necessity.

[0050] Example 1, in conjunction with Comparative Example 6, shows that Comparative Example 6 reduced the total reduction of area in cold drawing to 60%, with significantly lower elongation and reduction of area compared to Example 1, and a narrower interfacial carbon concentration gradient. This indicates that severe cold deformation with a total reduction of area ≥85% is a necessary condition for constructing sufficient high-density dislocation and interfacial defect channels, providing adequate diffusion paths for short-range carbon atom diffusion during subsequent low-temperature heat treatment. Insufficient deformation will significantly reduce the interfacial control effect, making it impossible to achieve the high plasticity target described in this application.

[0051] The above embodiments and comparative examples fully demonstrate the necessity and superiority of the synergistic technical solution of "alloy composition-rapid cooling-large deformation cold drawing-low temperature heat treatment" described in this application.

[0052] 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 preparing high-strength, high-ductility steel cord, characterized in that, Includes the following steps: S1. Obtain a cord steel billet, heat, roll, and cool the cord steel billet to obtain cord steel wire. The microstructure of the cord steel wire is pearlite with an average interlamellar spacing of 95-100 nm. By mass percentage, the composition of the cord steel billet is: C: 0.80%-0.85%, Si: 0.15%-0.30%, Mn: 0.40%-0.70%, Cr: 0.15%-0.30%, V: 0.05%-0.15%, N: 0.0080%-0.0120%, P≤0.012%, S≤0.008%, with the remainder being Fe and unavoidable impurities. S2. The cord steel wire is cold-drawn to obtain cord steel wire, wherein the total area shrinkage of the cold drawing is ≥85%; S3. The cord steel wire is heat-treated to obtain high-strength ductile steel cord, wherein the heat treatment holding temperature is 280-300℃.

2. The method for preparing a high-strength, high-ductility steel cord according to claim 1, characterized in that, In step S1, the heating temperature is 1020-1080℃.

3. The method for preparing a high-strength, high-ductility cord steel according to claim 1, characterized in that, In step S1, the cooling method is specifically as follows: cooling to 480-520℃ at a cooling rate of 8-12℃ / s and holding at that temperature for 10-30 seconds.

4. The method for preparing high-strength, high-ductility steel cord according to claim 1, characterized in that, In step S2, the number of cold drawing passes is 3-5.

5. The method for preparing a high-strength, high-ductility steel cord according to claim 1, characterized in that, In step S3, the heat treatment holding time is 30-90 minutes.

6. The method for preparing a high-strength, high-ductility steel cord according to claim 1, characterized in that, In step S3, when the temperature of the heat treatment is ≥280℃ and ≤290℃, the heat treatment holding time is ≥60min and ≤90min.

7. The method for preparing a high-strength, high-ductility cord steel according to claim 1, characterized in that, In step S3, when the temperature of the heat treatment is >290℃ and ≤300℃, the heat treatment holding time is ≥30min and <60min.

8. A high-strength, high-ductility steel cord, characterized in that, It is prepared by the method for preparing high-strength and ductile cord steel according to any one of claims 1-7.

9. A high-strength, ductile steel cord according to claim 8, characterized in that, The high-strength plastic cord steel has a tensile strength ≥2000MPa, a yield strength ≥1800MPa, an elongation ≥8.5%, and a reduction of area ≥40%.

10. A high-strength, ductile steel cord according to claim 8, characterized in that, The microstructure of the high-strength and ductile cord steel includes: pearlite and (Cr,V)CN composite precipitates. The average interlamellar spacing of the pearlite is ≤95nm. The pearlite is composed of cementite and ferrite, wherein carbon atoms diffuse from the cementite to the ferrite, forming a carbon concentration gradient transition region with a width of 50-100nm at the interface between the cementite and the ferrite. The average size of the (Cr,V)CN composite precipitates is ≤10nm, and the (Cr,V)CN composite precipitates are dispersed in the ferrite.

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