Low-yield-ratio high-strength hot-rolled strip steel resistant to deformation and cracking, and manufacturing method therefor

Abstract

The present invention relates to a hot-rolled strip steel resistant to deformation and cracking. The chemical components of the hot-rolled strip steel comprise, in percentages by weight: C: 0.05-0.14%, Si: 0.01-0.60%, Mn: 1.00-1.80%, P≤0.030%, S≤0.006%, Al: 0.01-0.30%, Ti: 0.08-0.18%, N≤0.006%, and the balance comprising Fe and other inevitable impurity elements. The chemical components of the hot-rolled strip steel further satisfies: the effective Ti content pTi≥0.06%; pTi=Ti-3.4N-3S; and the phase transition precipitation index n(α)=0.10-0.20%, and n(α)=C+Mn / 20-(Si+Al) / 15, wherein during calculation, the weight percentages of corresponding elements in the hot-rolled strip steel are substituted. Also disclosed in the present invention is a method for manufacturing the hot-rolled strip steel.

Description

A high-strength hot-rolled strip with low yield strength ratio and resistant to deformation cracking, and its manufacturing method. Technical Field

[0001] This invention relates to the field of low alloy steel technology, and in particular to a high-strength hot-rolled strip steel with a low yield strength ratio that is resistant to deformation and cracking, and its manufacturing method. Background Technology

[0002] Hot-rolled U-shaped steel is commonly used as a support component in coal mine roadway support structures. However, this method suffers from drawbacks such as low strength, heavy components, and difficulty in underground operations. To address these issues, hot-rolled sheet roll-formed profiles have been developed to replace U-shaped steel, achieving high strength and lightweight construction. Roll-formed profiles have complex cross-sections, requiring materials with excellent formability. Furthermore, roadway support materials must withstand long-term soil compression underground; therefore, the materials need excellent resistance to deformation and cracking to prevent localized fracture failure and collapse. Similar to the seismic and deformation-resistant design of building structures, the materials require a low yield strength ratio to withstand gradually increasing deformation and high toughness and plasticity to withstand deformation and crack propagation, ensuring the overall safety of the underground roadway support system. Although a low yield strength ratio is required, the high yield strength of over 600 MPa is still necessary to meet the lightweight requirements due to underground compressive strength demands.

[0003] However, existing high-strength steels with low yield strength ratios still have many shortcomings. For example:

[0004] Chinese patent application CN200710039741.0 discloses a steel plate with low-temperature toughness, high tensile strength, and low yield strength ratio. Through a combination of alloying element design and optimized TMCP process, a microstructure is formed consisting of a small amount of fine ferrite (10%–20%) and dispersed acicular ferrite, resolving the conflicting and difficult-to-reconcile issues of low-temperature impact toughness and low yield strength ratio in composition and process design. Controlling the yield strength ratio of the steel plate below 0.85 can improve the safety stability and earthquake resistance of steel structures. Based on its high heat input welding requirements… Its composition includes: C: 0.025%–0.055%, Si: ≤0.20%, Mn: 1.40%–1.75%, P ≤0.013%, S ≤0.002%, Cu: 0.25%–0.55%, Ni: 0.40%–0.80%, Mo: 0.10%–0.40%, Nb: 0.020%–0.040%, Als: 0.040%–0.060%, Ti: 0.007%–0.013%, N ≤0.0040%, Ca: 0.001%–0.005%, B ≤0.0003%. The high levels of Cu, Ni, and Mo in its alloy design will lead to higher production costs. Furthermore, based on the proposed design, the strength of the steel plate is below 553 MPa, which is a low strength grade and does not meet the requirements for coal mine roadway support materials.

[0005] Japanese invention JP2007071562 uses a simple steel composition, including: C: 0.07-0.18%, Si: 0.05-0.7%, Mn: 1.0-1.7%, sol.Al: 0.005-0.08%, and N: 0.001-0.008%, to achieve a high-strength steel with a yield strength ratio (YR) of less than 75%. The microstructure of this steel consists of polygonal ferrite phase with an area ratio of 65-85%, banded pearlite phase of 5-20%, and granular tempered martensite phase present at the ferrite grain boundaries of 3-15%. However, based on the proposed design, the yield strength achieved by this steel is 450-538 MPa, far less than 600 MPa.

[0006] Chinese patent application CN200510047979.9 discloses a hot-rolled duplex steel containing Nb and Ti, with the following composition: C 0.06–0.10%, Mn 0.4–1.2%, Si 0.015–0.45%, Nb 0.015–0.04%, Ti 0.005–0.03%, S≤0.003%, Al≤0.034%, P≤0.014%, and the balance being Fe. This steel exhibits a tensile strength of 500–610 MPa, an elongation of 20–30%, and a yield strength ratio ≤0.7. However, due to the low content of Mn, Nb, and Ti in the steel, effective strengthening design is not achieved, resulting in low strength.

[0007] Chinese patent application CN200710042357.6 discloses a high-strength, low-yield-strength-ratio steel. Through the combined design of alloying elements and a special TMCP+α / γ two-phase region normalizing + tempering process, a steel with a yield strength ≥650 MPa, tensile strength ≥780 MPa, and yield-strength ratio ≤0.80 was obtained. The microstructure of this steel consists of fine ferrite + tempered martensite (including some lower bainite) + a very small amount of tempered high-carbon martensite. The steel composition includes: C: 0.055–0.095%, Si ≤ 0.20%, Mn: 0.90–1.20%, P ≤ 0.013%, S ≤ 0.003%, Cr: 0.40–0.90%, Cu: 0.25–0.55%, Ni: 0.60–1.20%, Mo: 0.30–0.50%, Nb: 0.010–0.030%, Als: 0.040–0.060%, Ti: 0.008–0.020%, V: 0.015–0.055%, N ≤ 0.0040%, Ca: 0.001–0.005%, B: 0.0008–0.0016%. The material contains high levels of precious alloys such as Cr, Cu, Ni, and Mo, and its production process is extremely complex.

[0008] Chinese patent application CN201010599469.3 also discloses an 800MPa grade low yield strength ratio structural steel plate, the composition of which includes C: 0.045-0.075%, Si: 0.30-0.55%, Mn: 1.55-1.95%, P≤0.01%, S≤0.0025%, Alt: 0.012-0.035%, Cr: 0.15-0.25%, Mo: 0.15-0.3%, Cu: 0.2-0.4%, Ni: 0.2-0.4%, Nb: 0.008-0.04%, V: 0.008-0.04%, Ti: 0.008-0.03%, and B: 0.0008-0.0015%. The TMCP+tempering heat treatment is used to achieve a yield strength ≥550MPa, tensile strength ≥800MPa, and yield-to-tensile ratio <0.70.

[0009] Chinese patent application CN200810200100.3 discloses a high-strength steel plate with a low yield strength ratio, achieving a strength of 1200–1500 MPa. Its chemical composition includes: C: 0.15–0.20 wt%, Si: 1.0–2.0 wt%, Mn: 1.8–2.0 wt%, Al ≤ 0.036 wt%, V: 0.05–0.1 wt%, P ≤ 0.01 wt%, S ≤ 0.005 wt%, and Cr: 0.8–1.0 wt%. The steel employs a high-C, Mn, and Cr alloy design, with a microstructure of bainite or bainite + martensite. It is a low-temperature transformation strengthened steel, but exhibits relatively low plasticity and formability.

[0010] Chinese patent application CN201210470414.1 also discloses a high-strength multiphase steel plate with a yield strength >700MPa, tensile strength >1000MPa, elongation A80 >10%, and yield strength ratio <0.75. The chemical composition includes: C: 0.175–0.215%, Si: 0.15–0.55%, Mn: 1.60–2.00%, Al: 0.015–0.040%, Ti: 0.005–0.020%, Nb: 0.01–0.03%, N: ≤0.006%. A segmented manufacturing process is used to obtain a multiphase microstructure of bainite + martensite (possibly containing a small amount of ferrite). In summary, this also achieves high-strength, low-yield-strength-ratio steel through high carbon content and a two-stage cold + low-temperature coiling process, but due to the low plasticity of the bainite / martensite structure, it is difficult to meet the aforementioned processing objectives.

[0011] In summary, the existing patents disclose that all low yield strength ratio hot-rolled steels primarily employ microstructural strengthening as their strengthening mechanism. Due to welding and application requirements, some designs with simple compositions cannot achieve the high yield strength requirement of over 600 MPa. Others utilize high-content Cu, Cr, Ni, Mo, V, Nb compositions and complex TMCP + tempering processes to achieve microstructural strengthening steels with a fine ferrite + tempered martensite / lower bainite microstructure, achieving tensile strengths of 800 MPa (yield strength above 550 MPa). However, these steels are costly and have complex production processes. Some patents also employ higher carbon content (C content ≥ 0.15%) with low-temperature coiling or two-stage cooling to achieve bainite / martensite or its multiphase microstructure, resulting in ultra-high strength low yield strength ratio steels. However, due to their lower plasticity, these steels are difficult to meet target processing requirements. Summary of the Invention

[0012] The purpose of this invention is to provide a high-strength hot-rolled strip steel with a low yield strength ratio and resistant to deformation cracking, and a method for manufacturing the same. This hot-rolled strip steel possesses high strength, a low yield strength ratio, and high toughness and plasticity. The preferred properties of the hot-rolled strip steel are: yield strength ≥ 600 MPa, tensile strength ≥ 700 MPa, yield strength ratio ≤ 0.85, and elongation ≥ 16.0%. It meets the requirements for 180° cold bending with a bending diameter twice the plate thickness, and after a 90° bend, it exhibits crack resistance without crack propagation upon flattening. Considering factors such as manufacturing cost, the yield strength of the hot-rolled strip steel in this invention is typically below 800 MPa, the tensile strength is typically below 1050 MPa, the yield strength ratio is between 0.60 and 0.85 (e.g., 0.63-0.85), and the elongation is typically below 22.0% (e.g., below 21.1%).

[0013] The steel plate of this invention is suitable for producing roll-formed profiles with complex cross-sections to replace hot-rolled U-shaped steel; it can meet the lightweight requirements of arched support components for coal mine roadways, can withstand the continuous increase of mine pressure in underground coal mine roadways and the deformation requirements of arched support structures, can resist deformation cracking and crack propagation, and delay local node fracture.

[0014] To achieve the above objectives, the technical solution of the present invention is as follows:

[0015] This invention utilizes a suitable C-Mn-Ti-based alloy composition design to achieve synergistic control of TiC interphase precipitation and low-temperature phase transformation microstructure formation in the steel matrix. Based on traditional low yield strength ratio alloy steel, the ferrite is strengthened by interphase precipitation of TiC particles with an average size of less than 10 nm, effectively improving the yield strength. Based on traditional high-strength steel, a lower yield strength ratio and good resistance to deformation cracking are achieved through ferrite + low-temperature phase transformation microstructure.

[0016] Specifically, the hot-rolled strip steel resistant to deformation cracking described in this invention comprises the following chemical composition by weight percentage: C 0.05-0.14%, Si 0.01-0.60%, preferably 0.10-0.60%, Mn 1.00-1.80%, P≤0.030%, S≤0.006%, Al 0.01-0.30%, Ti 0.08-0.18%, N≤0.006%, with the balance including Fe and unavoidable impurity elements. Furthermore, the chemical composition of the hot-rolled strip steel also satisfies the following:

[0017] Effective Ti content p Ti ≥0.06%; p Ti =Ti-3.4N-3S;

[0018] The phase change precipitation index n(α) = 0.10~0.20%, n(α) = C + Mn / 20 - (Si + Al) / 15;

[0019] When calculating, substitute the weight percentage of the corresponding element in the hot-rolled strip steel.

[0020] It should be noted that the calculated values ​​of the formulas involved in this invention are allowed to have conventional errors due to measurement accuracy or calculation rounding.

[0021] Preferably, the hot-rolled strip steel further comprises one or more elements selected from the following: Nb 0.003-0.050%, V 0.005-0.100%, Cr 0.01-0.60%, Mo 0.01-0.30%, Ni 0.01-0.30%, Cu 0.01-0.50%, B 0.0003-0.0020%, and Cr+Mo+Ni+Cu≤0.70%, where the element symbols represent the weight percentage of the corresponding element in the hot-rolled strip steel.

[0022] Preferably, the hot-rolled strip steel further comprises one or more elements selected from Ca 0.0004-0.0050%, Mg 0.0004-0.0050%, and rare earth Re 0.0004-0.0050%.

[0023] In this invention, rare earth (Re) refers to one or more of the lanthanides, scandium, or yttrium. Preferably, the rare earth is an industrial mixed rare earth, such as a cerium-rich mixed rare earth or a lanthanum-rich mixed rare earth, wherein the main components are Ce and / or La.

[0024] Preferably, the chemical composition of the hot-rolled strip steel contains the balance of Fe and unavoidable impurity elements.

[0025] The microstructure of the hot-rolled strip steel described in this invention is: 30-70% polygonal ferrite in area + 25-65% acicular ferrite, bainite, lower bainite or martensite, or one or more of these + 5-35% modified pearlite, which is unavoidable in continuous phase transformation.

[0026] Preferably, in the microstructure of the hot-rolled strip steel, nano-sized Ti carbide particles are distributed in the polygonal ferrite microstructure in the form of interphase precipitation, and more than 80% of the Ti carbide particles have a diameter of less than 10 nm.

[0027] Unless otherwise specified, the diameter of the carbide particles in this invention refers to the diameter of the equivalent circle.

[0028] The hot-rolled strip steel of the present invention has a yield strength ≥600MPa, tensile strength ≥700MPa, yield strength ratio ≤0.85, elongation ≥16.0%, and preferably can also meet the requirements of 180° cold bending with a bending diameter of 2 times the plate thickness, and can meet the anti-cracking performance of flattening without crack propagation after 90° bending.

[0029] The hot-rolled strip steel of the present invention has the following performance characteristics: it has high strength, which can meet the lightweight requirements of the arched components of roadway support; it has a low yield strength ratio, which can withstand the continuous increase of mine pressure in underground roadways and the deformation requirements of the arched support structure; it has high toughness and plasticity, which can resist deformation cracking and crack propagation, and delay local node fracture.

[0030] The composition design of the hot-rolled strip steel in this invention is explained as follows:

[0031] Carbon (C) is an effective strengthening element in steel. Besides solid solution strengthening, it can also form nanoscale second-phase precipitates with microalloying elements such as Ti and Nb, playing a role in precipitation strengthening and microstructure refinement. Simultaneously, in conjunction with processing, C can increase the content of low-temperature transformation structures such as bainite or martensite, as well as hardness, thereby significantly improving the tensile strength of the material and achieving a lower yield strength ratio. However, excessive C can lead to excessively high tensile strength, reducing the material's plasticity and formability, and also decreasing its toughness. Therefore, the C content in this invention is 0.05–0.14%.

[0032] Si is a commonly used deoxidizing element in steel and also has a solid solution strengthening effect on steel. However, a high Si content will affect the surface quality of the strip steel and reduce the toughness of the material. Therefore, the Si content in this invention is 0.01 to 0.60%, preferably 0.10 to 0.60%.

[0033] Mn is an important strengthening and toughening element in steel, playing a role in solid solution strengthening. It also lowers the transformation temperature of supercooled austenite and the ferrite transformation temperature, which is beneficial for microstructure refinement and improves the strength and toughness of the material. Simultaneously, it inhibits the high-temperature induced precipitation of Ti in austenite, facilitating the precipitation of TiC at lower temperatures and enhancing its strengthening ability. Too low a Mn content leads to an excessively high ferrite proportion and a low content of hard phases such as bainite after two-stage cooling, resulting in a high yield strength ratio. However, excessive Mn content inhibits ferrite transformation, making interphase precipitation of TiC difficult and reducing the precipitation strengthening effect. Furthermore, low ferrite content also reduces the material's plasticity and cold formability. Therefore, the Mn content in this invention is 1.00–1.80%.

[0034] Phosphorus (P) is a harmful impurity element in steel, and it tends to segregate in the center of the thickness during continuous casting of steel billets. Simultaneously, P easily aggregates at grain boundaries, reducing the grain boundary binding energy and thus decreasing the toughness and plasticity of the steel. Therefore, the P content in steel should be minimized; this invention requires a P content ≤ 0.030%.

[0035] Sulfur (S) is a common harmful impurity element in steel, negatively impacting low-temperature toughness, weldability, and cold formability. It can combine with manganese (Mn) to form ductile MnS inclusions, which, during rolling, form elongated inclusions, reducing the steel's formability. Furthermore, S can react with titanium (Ti) at high temperatures to form TiS or Ti₄S₂C₂, reducing the content of effective Ti that can combine with carbon (C). Therefore, this invention requires an S content ≤ 0.006%.

[0036] Al is a very effective deoxidizing element. It also helps refine grains, improves the strength and toughness of steel, and promotes ferrite transformation. However, higher Al content is detrimental to continuous casting, easily clogging the nozzle. Therefore, the Al content in this invention is 0.01–0.30%.

[0037] Ti is a strong carbonitride forming element that can precipitate as nano-sized TiC second-phase particles, significantly improving the strength of materials. Therefore, it is added as an important strengthening element in this invention. During two-stage cooling, the solubility of TiC in steel decreases significantly during the isothermal transformation from austenite to ferrite. Consequently, TiC nucleates and precipitates at the austenite-ferrite phase interface. As the ferrite grain boundaries move forward, this forms row-like or layered nano-precipitates, greatly increasing the strength of the ferrite matrix. However, when the content is too high, the precipitation strengthening effect of Ti gradually weakens, and its tendency to coarsen significantly reduces the low-temperature toughness of the steel. Therefore, the Ti content in this invention is 0.08–0.18%.

[0038] Nitrogen (N) is an impurity element in steel. When N combines with Ti in molten steel, it can form large-sized TiN inclusions. This reduces the effective content of Ti and also significantly impairs the toughness of the steel. Therefore, the N content should be controlled as low as possible. In this invention, the N content is ≤0.006%.

[0039] Furthermore, this invention requires the steel to have an appropriate effective Ti content. The effective Ti content is related to the impurity elements in the steel; N, S, and C can all combine with Ti to form Ti compounds such as TiN, TiS, or Ti4C2S2 and TiC. However, before forming TiC, Ti preferentially combines with N and S in the steel. These Ti compounds precipitate at high temperatures, and the particles tend to grow quite large, typically with a particle size greater than 50 nm. According to Orowan strengthening theory, precipitated particles of this size contribute very little to the strengthening of the steel. Therefore, the remaining Ti content precipitated in the form of TiN and TiS is considered the effective Ti content that contributes to strengthening. In this invention, the remaining effective Ti content p is calculated using the following formula. Ti p Ti =Ti-3.4N-3S, when calculating, substitute the weight percentage of the corresponding element in the hot-rolled strip steel. Therefore, in order to ensure the desired precipitation strengthening contribution, this invention also requires the effective Ti content p in the steel. Ti ≥0.060%.

[0040] Furthermore, this invention limits the phase transformation precipitation index n(α): 0.10% ≤ n(α) ≤ 0.20%, where n(α) = C + Mn / 20 - (Si + Al) / 15, and the weight percentage of the corresponding element in the hot-rolled strip is substituted during calculation. The phase transformation precipitation index is an important parameter for coordinating TiC interphase precipitation and ferrite phase transformation. Compared with traditional hot-rolled dual-phase steel, this invention achieves ferrite strengthening through the TiC interphase precipitation process, thereby realizing a high-yield-strength steel with a low yield-to-tensile ratio using a very economical C-Mn-Ti composition system.

[0041] During the cooling process of steel, as austenite (γ) transforms into ferrite (α), carbon atoms, being readily diffusible, rapidly accumulate at the phase boundaries due to their significantly lower solid solubility in the α phase compared to γ. This leads to a substantial increase in the solid solubility product of Ti and C at the phase boundaries. Simultaneously, the solubility of TiC in ferrite decreases significantly, resulting in a high precipitation driving force for TiC, which then nucleates along the two-phase boundaries. As the γ→α transformation occurs, ferrite grows, and the phase boundaries migrate. The TiC particles that initially precipitated along these boundaries are now distributed in rows or layers within the ferrite matrix, while new nucleation begins at the migrated γ / α phase boundaries. This is the interphase precipitation process of TiC. The resulting densely packed TiC precipitates in ferrite are small in size (average equivalent circle diameter ≤ 10 nm), dense, and closely spaced, providing extremely strong strengthening to the soft ferrite phase. Therefore, the synergistic effect of TiC interphase precipitation and ferrite phase transformation is crucial in this invention.

[0042] In this invention, the phase transformation precipitation index n(α) directly affects the properties of the steel. When n(α) is below 0.10%, the ferrite phase transformation temperature of the material becomes too high, the transformation rate becomes too fast, and ultimately, an excessive amount of ferrite phase change is obtained. With an excessively high ferrite phase transformation temperature, the undercooling of TiC particle precipitation is low, resulting in a low nucleation density. Simultaneously, after TiC precipitation, due to the high temperature, it tends to coarsen, resulting in fewer TiC particles, larger particle size, and larger spacing. According to the Orowan strengthening mechanism, the strengthening effect of these particles is weak. Furthermore, an excessively fast ferrite phase transformation rate also leads to Fe3C precipitation, reducing the C atom concentration at grain boundaries, affecting the bonding between Ti and C, and significantly reducing the number of TiC precipitated particles. Additionally, if the ferrite transformation rate is too high, such as exceeding 70%, it also leads to a low proportion of hard phase structure, resulting in a yield strength ratio exceeding 0.85.

[0043] When n(α) is higher than 0.20%, the stability of austenite increases, making ferrite phase transformation difficult during laminar cooling of the strip. Without ferrite phase transformation, the conditions for TiC interphase precipitation are lost. Therefore, excessively high n(α) affects TiC interphase precipitation, leading to a decrease in yield strength. Simultaneously, excessively high n(α) reduces the amount of ferrite phase change, causing the material's microstructure to transform more into hard phases such as bainite / martensite. This results in excessively high tensile strength, reduced elongation, and negatively impacts processing formability and resistance to deformation.

[0044] Therefore, the present invention limits 0.10% ≤ n(α) ≤ 0.20%, which is the key to obtaining a suitable ferrite structure ratio and the interphase precipitation of Ti carbides. It is also the key to the ferrite being strengthened by the interphase precipitated Ti carbides to obtain a relatively high yield strength.

[0045] Elements such as Nb and V are also strong carbide-forming elements, and can form carbide particles such as NbC and VC, resulting in precipitation strengthening. However, the cost of Nb and V is much higher than that of Ti, so their economic efficiency in improving strength is not as good as that of Ti. Therefore, although Nb and V are not preferred precipitation strengthening elements in this invention, they can still be added to the steel in appropriate amounts.

[0046] Elements such as Cu, Cr, Ni, Mo, and B are commonly used strengthening alloying elements in steel. They increase hardenability, inhibit ferrite transformation, refine microstructure, and enhance the contribution of TiC precipitation strengthening. Adding small amounts is beneficial to the strength and toughness of steel. However, Cu, Cr, Ni, and Mo are all precious metal elements, and excessive amounts are detrimental to the economics of steel production. Furthermore, Cu, Cr, Ni, and Mo all inhibit ferrite transformation; therefore, the total amount of these elements must be limited to avoid insufficient ferrite content, which would reduce the material's formability. In addition, excessive B can increase grain boundary brittleness, which is detrimental to the material's toughness.

[0047] Therefore, the present invention specifies that the steel may contain one or more of Nb, V, Cr, Mo, Ni, Cu, and B, with the following content: Nb: 0.003–0.05%, V: 0.005–0.10%, Cr: 0.01–0.60%, Mo: 0.01–0.30%, Ni: 0.01–0.30%, Cu: 0.01–0.50%, B: 0.0003–0.0020%, and Cr+Mo+Ni+Cu ≤ 0.70%. In calculations, the weight percentage of the corresponding element in the hot-rolled strip steel is substituted. This can further improve the strength and toughness of the steel of the present invention.

[0048] Ca can form spherically dispersed CaS with S, thereby improving the distribution of MnS inclusions in steel and preventing soft MnS inclusions from elongating during rolling and causing delamination during steel shearing. This invention limits the Ca content to 0.0004-0.0050%.

[0049] Mg can combine with O to form finer MgO, which can serve as nucleation sites for other inclusions, thereby refining the size of inclusions, improving their dispersion, and enhancing the toughness of the material. This invention limits the Mg content to 0.0004–0.0050%.

[0050] Rare earth element retinylamine (re) also improves corrosion resistance and increases material toughness. However, the yield of retinylamine during steelmaking is difficult to control, and excessive addition will reduce the economics of steel. Therefore, this invention limits the retinylamine content to 0.0004–0.0050%.

[0051] Ca, Mg, Re, etc. are all elements that improve the quality of molten steel and reduce harmful inclusions in the steel during the steelmaking process. Therefore, adding one or more of them can be beneficial to improving the internal quality and toughness of the steel.

[0052] The present invention also provides a method for manufacturing the aforementioned deformation-resistant, low-yield-strength, high-strength hot-rolled strip steel and its manufacturing method, comprising the following steps:

[0053] 1) Smelting and casting

[0054] Based on the above-mentioned chemical composition of hot-rolled strip steel, it is smelted and cast into slabs;

[0055] 2) Slab heating

[0056] The slab is heated in the heating furnace of the hot rolling production line or in the homogenizing furnace of the thin slab continuous casting and rolling production line; the outlet temperature of the heating furnace in the hot rolling production line is 1230-1280℃, and the heating time above 1230℃ is 30-60 minutes; the outlet temperature of the homogenizing furnace in the thin slab continuous casting and rolling production line is 1150-1230℃.

[0057] 3) Rolling

[0058] After the slab exits the heating furnace or soaking furnace, it undergoes rough rolling. After rough rolling, it enters the finishing mill and uses a multi-stand continuous rolling process to obtain strip steel, controlling the final rolling temperature to 800-880℃.

[0059] 4) Laminar flow cooling

[0060] After exiting the rolling mill, the strip undergoes laminar flow cooling in a segmented manner. The first segment uses water cooling with a controlled cooling rate ≥100℃ / s. Then, the strip enters an air-cooling section. The initial temperature of the air-cooling section is T±20℃, where T=700-n(α) / 0.15×50℃, and the air-cooling time is t±1s, where t=[n(α) / 0.15+p Ti [0.2]×10s; After the air cooling section ends, the second water cooling section begins, where the strip is cooled to below 200℃ at a cooling rate of ≥60℃ / s before being coiled.

[0061] Unless otherwise specified, the cooling rate of the strip in this invention refers to the average cooling rate.

[0062] The process of this invention is mainly designed to better control the precipitation of Ti carbides and the transformation of the microstructure in steel with the above-mentioned composition, thereby improving material properties. The requirement for the billet to be heated in the furnace at above 1230°C for 30–60 minutes ensures sufficient Ti solution dissolution. The final rolling temperature is controlled at 800–880°C, a lower temperature than conventional steel rolling processes. This is to increase the rolling deformation energy and promote the transformation of fine-grained ferrite during the cooling process of the strip layer. This is crucial for promoting the precipitation of Ti carbide particles to improve strength and increasing ferrite content to improve plasticity in the steel of this invention.

[0063] During the laminar flow cooling stage, the precipitation and microstructure of the invented steel are precisely controlled through a segmented cooling process. The cooling rate of the first stage of laminar flow cooling is ≥100℃ / s. Rapid cooling can significantly increase the undercooling degree of phase transformation, increase the number of phase deformation nuclei, and refine the polymorphic ferrite.

[0064] After the first stage of laminar flow cooling, the strip enters an air-cooled state. Under near-isothermal or low cooling rate conditions, the material undergoes a ferrite phase transformation and interphase precipitation of Ti carbide particles. The phase transformation rate and phase change of ferrite, as well as the precipitation rate and amount of Ti carbide particles, are closely related to the air-cooling temperature and time. This invention is based on the phase transformation precipitation index n(α) related to the composition of C, Mn, Si, Al, and Ti, and the effective Ti content p. Ti By determining the appropriate air-cooling temperature and time, Ti carbide particles were precipitated in the ferrite as interphase precipitation during the ferrite transformation of the strip. Over 80% of these precipitated particles had a diameter (equivalent circular diameter) below 10 nm and exhibited a typical interphase precipitation pattern (row-like or layered distribution). This resulted in a strengthening contribution of 100–200 MPa through the Orowan strengthening mechanism, achieving hardening of the soft ferrite phase and thus effectively improving the material's yield strength.

[0065] Following the air-cooling section, the strip enters the second stage of cooling, with a cooling rate ≥60℃ / s. Combined with the phase transformation precipitation index n(α), this transforms the untransformed austenite into a corresponding proportion of low-temperature phase transformation structures. These low-temperature structures are complex, and depending on the carbon content and alloying elements, they can form one or more of acicular ferrite, bainite, lower bainite, or martensite, such as composite structures. These hard phase structures have very high strength, significantly improving the tensile strength of the material, and form a soft-hard phase composite structure with the relatively soft ferrite phase, thereby achieving a relatively low yield strength ratio.

[0066] The main innovation of this invention compared with the prior art is as follows:

[0067] Existing high-strength steels with low yield strength ratios typically have yield strengths below 550 MPa. To achieve higher strength levels, a complex TMCP (Transformer Metal Processing) and tempering process is usually employed, using a composition system of Cu, Cr, Ni, Mo, V, and Nb. Because strength enhancement is primarily achieved through microstructural strengthening, the alloy cost is high and the production process is complex. Alternatively, a higher carbon content (C content ≥ 0.15%) combined with a low-temperature coiling process can be used to form a bainitic / martensite or multiphase low-temperature microstructure, resulting in high-strength, low-yield-strength steel. However, these low-temperature microstructure steels have lower plasticity, making it difficult to meet the processing requirements of materials such as tunnel support.

[0068] This invention combines the interphase precipitation behavior of TiC with the two-phase or multiphase microstructure control of traditional low yield strength ratio steel by designing appropriate composition and phase transformation precipitation index n(α). It achieves the design of high yield strength and low yield strength ratio high-strength steel with a very economical C-Mn-Ti composition system, and obtains hot-rolled strip steel that can meet the processing requirements of complex cross-section profiles such as coal mine roadway support and has excellent properties such as high compressive strength and fracture resistance.

[0069] Traditional C-Mn steels are produced as dual-phase steels with very low yield strength ratios through a segmented cooling process. During cooling, the steel is held in the two-phase region at a temperature of 600-700℃, where austenite transforms into ferrite (γ→α), forming a certain amount of polygonal ferrite. The remaining untransformed austenite forms higher-strength low-temperature structures such as bainite and martensite during the second cooling stage. However, the polygonal ferrite formed in this way has very low strength, resulting in a low yield strength ratio, but also a very low yield strength.

[0070] This invention utilizes a rational composition control and a special technique and process for the interphase precipitation of Ti carbides to effectively enhance the strength of ferrite, thereby achieving a high yield strength and low yield ratio steel. This invention controls the content of the main elements C, Mn, and Ti, and establishes an effective Ti content (p0.05). Ti=Ti-3.4N-3S) and the control requirements for the phase transformation precipitation index (n(α)=C+Mn / 20-(Si+Al) / 15), p Ti With a concentration of ≥0.06% and 0.10%≤n(α)≤0.2%, it is possible to achieve simultaneous transformation of ferrite phase transformation and interphase precipitation of Ti carbides.

[0071] During the γ→α transformation in the two-phase temperature range, carbon, being a readily diffusible atom, accumulates at the phase boundaries, leading to a significant increase in the solid solubility of [Ti][C] at these boundaries. Simultaneously, the solubility of [Ti][C] decreases significantly after austenite transforms into ferrite, thus significantly increasing the driving force for TiC precipitation. Therefore, Ti carbides have sufficient thermodynamic conditions to precipitate and nucleate along the two-phase boundaries. With the γ→α transformation, ferrite grows, the phase boundaries shift forward, and the TiC particles that previously precipitated along the phase boundaries are distributed in rows or layers within the ferrite matrix, while new nucleation begins at the γ / α phase boundaries. As the phase transformation completes, TiC particles precipitate from the ferrite as interphase precipitates. These nano-sized (average equivalent circle diameter ≤10nm) TiC precipitates, formed by interphase precipitation, contribute extremely strongly to the strengthening of the soft ferrite phase (100-200MPa) due to their small size, high density, and small spacing, thereby achieving ferrite strengthening and significantly improving the yield strength of the material.

[0072] This invention designs a suitable material microstructure to give the material the desired properties. Combined with the above-mentioned p... Ti Ideal process parameters are determined using parameters such as n(α), for example, the optimal starting temperature of the air-cooling section is T±20℃, where T=700-n(α) / 0.15×50℃; the air-cooling time is t±1s, where t=[n(α) / 0.15+p Ti [0.2]×10s, further optimized control of ferrite phase change and Ti carbide precipitation was obtained, resulting in a suitable microstructure: polygonal ferrite (area percentage 30-70%), and a multiphase structure composed of one or more of acicular ferrite, bainite, lower bainite, or martensite (area percentage 25-65%). During the continuous phase transformation, non-equilibrium pearlite transformation inevitably occurs to form modified pearlite (area percentage 5-35%). Preferably, the polygonal ferrite contains interphase precipitates of nano-sized Ti carbides, with more than 80% of the Ti carbide particles having a diameter ≤10nm, which has a significant strengthening effect on ferrite. Through the strengthened ferrite and a suitable proportion of low-temperature structure (acicular ferrite, bainite, lower bainite, or martensite, or a composite structure of these structures), the material still has a relatively low yield strength ratio, but the yield strength is significantly improved compared to traditional dual-phase steel.

[0073] Therefore, this invention utilizes a rational design of the C, Mn, and Ti content, combined with p Ti The requirements of n(α) control the interphase precipitation behavior of Ti carbides, thereby achieving an ideal combination of precipitation and microstructure. This results in a material with a low yield strength ratio of ≤0.85, a high yield strength of ≥600MPa, and good toughness and plasticity. The elongation of the steel is ≥16%. Preferably, it also meets the requirements of 180° cold bending with a bending diameter of 2 times the plate thickness, and can be flattened without cracking after a 90° bend. This satisfies the material requirements of this invention for lightweight coal mine roadway support with complex cross-section roll forming, which has high compressive strength and fracture resistance. Attached Figure Description

[0074] Figure 1 is a micrograph of the hot-rolled strip steel of Embodiment 1 of the present invention;

[0075] Figure 2 is a micrograph of the hot-rolled strip steel of Example 3 of the present invention;

[0076] Figure 3 is a micrograph of the hot-rolled strip steel of Example 4 of the present invention;

[0077] Figure 4 is a photograph of the morphology of interphase precipitates in the hot-rolled strip steel of Example 1 of the present invention.

[0078] Figure 5 is a photograph of the hot-rolled strip steel of Embodiment 1 of the present invention after cold bending at 180° with D=2a (a is the plate thickness);

[0079] Figure 6 shows photographs of hot-rolled strip steel of Embodiment 1 of the present invention after being cold-bent at 90° and then flattened, with D=0.5a, D=1a, and D=2a (a being the plate thickness). Detailed Implementation

[0080] The present invention will be further described below with reference to the embodiments and accompanying drawings.

[0081] The chemical composition of the hot-rolled strip steel of the present invention is shown in Tables 1 and 2, the manufacturing process parameters are shown in Table 3, and the performance parameters of the hot-rolled strip steel of the embodiments are shown in Table 4.

[0082] The strip steel in the embodiment is manufactured as follows:

[0083] Steel billets were obtained by converter smelting, refining, and continuous casting according to the chemical composition shown in Table 1. The billets were heated in a heating furnace at a temperature of 1230–1280℃, and held at 1230℃ for 30–60 minutes before hot rolling and controlled cooling. Hot rolling was carried out in two stages. After rough rolling, the steel entered the finishing mill, which was a 7-stand continuous rolling mill with a final rolling temperature of 800–880℃, rolling to a thickness of 16 mm to 3 mm. After exiting the mill, the strip underwent laminar flow cooling, which was carried out in stages. After the first stage of cooling, the strip entered the air cooling section, and then entered the second stage of cooling to below 200℃. The specific process parameters are shown in Table 3.

[0084] As can be seen from Examples 1 to 47, when the steel composition meets the design requirements and the rolling and cooling are controlled according to the above process requirements, products that meet the performance requirements are obtained: yield strength ≥ 600 MPa, tensile strength ≥ 700 MPa, yield ratio ≤ 0.85, and elongation ≥ 16% (yield strength, tensile strength, and elongation are tested according to GB / T 228.1-2021 "Metallic materials, tensile testing - Part 1: Test at room temperature", where elongation is the elongation after fracture A5, which refers to the elongation after fracture measured using a proportional gauge length specimen with a proportionality coefficient k = 5.65). All the steels in all embodiments can also meet the requirements of 180° cold bending with a bending diameter of 2 times the plate thickness (tested according to GB / T 232-2024 "Metallic Materials Bending Test Method", no cracking), and after bending at 90° and then flattening in the reverse direction, no cracking or expansion is observed on the surface of the strip (the strip sample is bent at 90° according to GB / T 232-2024 "Metallic Materials Bending Test Method", and then the bent strip is flattened in the reverse direction, and no cracking or expansion is observed on the surface of the strip).

[0085] The typical microstructure of the hot-rolled strip steel of this invention is shown in Figures 1 to 4. After mechanical polishing and etching with 4% nitric acid alcohol, the cross-section of the steel plate was observed under a scanning electron microscope (SEM), and the area ratio of each phase was statistically analyzed using image analysis software. As shown in the figures, the material has a multiphase structure, with the microstructure conforming to a mixture of 30-70% polygonal ferrite, 25-65% acicular ferrite, bainite, lower bainite, or martensite, and 5-35% modified pearlite. Furthermore, transmission electron microscopy (TEM) revealed that nano-sized Ti carbide particles precipitated in the polygonal ferrite microstructure as interphase precipitation, with over 80% of the Ti carbide particles having a diameter of less than 10 nm and exhibiting a row-like distribution.

[0086] A cold bending test of 180° with D=2a (a is the plate thickness) with sheared edge was conducted on the strip of Example 1. As shown in Figure 5, the strip of Example 1 of the present invention did not crack, and the edge shear hardening layer did not have crack propagation, showing good formability and crack resistance.

[0087] The resistance to deformation cracking after bending the strip steel of Example 1 was further evaluated. As shown in Figure 6, the typical morphology of the steel of Example 1 after being cold-bent at 90° with D=0.5a, D=1a, and D=2a (a is the plate thickness) and then flattened is shown. The steel of Example 1 of the present invention exhibits good resistance to crack propagation.

[0088] Comparative Examples 1-4 show insufficient yield strength due to low effective Ti.

[0089] Comparative Example 5 has a low C content, resulting in a small phase transformation precipitation index and a low effective Ti content. This leads to an excessively large ferrite phase change, insufficient precipitation, or an insufficient proportion of small-sized particles, resulting in low yield strength, low tensile strength, and a yield-to-tensile ratio higher than 0.85.

[0090] In contrast, the low Mn content in Comparative Example 6 resulted in a smaller phase transformation precipitation index, leading to an excessively large ferrite phase change and a low ratio of bainite / martensite hard phases, resulting in a higher yield strength ratio.

[0091] In Comparative Examples 7-9, the high C or Mn content led to a large phase transformation precipitation index, which made ferrite transformation difficult and resulted in a low ferrite content. This suppressed the precipitation of Ti carbides, resulting in a low yield strength. At the same time, it increased the proportion of bainite / martensite hard phases, leading to high tensile strength, insufficient elongation, and a decrease in the material's formability. Cracking occurred after flattening following bending.

Claims

1. A hot-rolled strip steel resistant to deformation cracking, comprising the following chemical composition by weight percentage: C 0.050–0.140%, Si 0.01–0.60%, preferably 0.10–0.60%, Mn 1.00–1.80%, P ≤0.030%, S ≤0.006%, Al 0.010–0.300%, Ti 0.080–0.180%, N ≤0.006%, with the balance including Fe and unavoidable impurity elements, and the chemical composition of the hot-rolled strip steel further satisfying: Effective Ti content p Ti ≥0.06%; p Ti =Ti-3.4N-3S; The phase change precipitation index n(α) = 0.10~0.20%, n(α) = C + Mn / 20 - (Si + Al) / 15; The element symbols in the formula represent the weight percentage of the corresponding element in the hot-rolled strip steel.

2. The hot-rolled strip steel resistant to deformation cracking as described in claim 1, characterized in that, The hot-rolled strip steel further comprises one or more elements selected from Nb, V, Cr, Mo, Ni, Cu, and B, wherein the content of each element satisfies the following: Nb 0.003~0.050%, V 0.005~0.100%, Cr 0.01~0.60%, Mo 0.01~0.30%, Ni 0.01~0.30%, Cu 0.01~0.50%, B 0.0003~0.0020%, and Cr+Mo+Ni+Cu≤0.70%, where the element symbol represents the weight percentage of the corresponding element in the hot-rolled strip steel.

3. The hot-rolled strip steel resistant to deformation cracking as described in claim 1 or 2, characterized in that, The hot-rolled strip steel also contains one or more elements selected from: Ca 0.0004-0.0050%, Mg 0.0004-0.0050%, and rare earth Re 0.0004-0.0050%.

4. The hot-rolled strip steel resistant to deformation cracking as described in claim 1, 2, or 3, characterized in that, The chemical composition of the hot-rolled strip steel contains the remainder Fe and unavoidable impurity elements.

5. The hot-rolled strip steel resistant to deformation cracking as described in claim 1, 2, 3, or 4, characterized in that, The microstructure of the hot-rolled strip is: 30-70% polygonal ferrite in area + 25-65% acicular ferrite, bainite, lower bainite or martensite or one or more of these in area + 5-35% modified pearlite that is unavoidable in continuous phase transformation; preferably, nano-sized Ti carbide particles precipitated in the form of interphase precipitation are distributed in the polygonal ferrite microstructure, and more than 80% of the Ti carbide particles have a diameter of less than 10 nm.

6. The hot-rolled strip steel resistant to deformation cracking as described in claim 1, 2, 3, 4, or 5, characterized in that, The hot-rolled strip has a yield strength ≥600MPa, tensile strength ≥700MPa, yield strength ratio ≤0.85, and elongation ≥16.0%. Preferably, it also meets the requirements of 180° cold bending with a bending diameter of 2 times the plate thickness, and can meet the anti-cracking performance of flattening without crack propagation after 90° bending.

7. A method for manufacturing hot-rolled strip steel resistant to deformation cracking as described in any one of claims 1 to 6, characterized in that, The method includes the following steps: 1) Smelting and casting The hot-rolled strip steel is smelted based on the chemical composition of any one of claims 1-4 and then cast into slabs. 2) Slab heating The slab is heated in the heating furnace of the hot rolling production line or homogenized in the soaking furnace of the thin slab continuous casting and rolling production line; wherein... The furnace outlet temperature of the heating furnace in the hot rolling production line is 1230-1280℃, and the heating time above 1230℃ is 30-60 minutes; the furnace outlet temperature of the homogenization furnace in the thin slab continuous casting and rolling production line is 1150-1230℃. 3) Rolling After the slab exits the heating furnace or soaking furnace, it undergoes rough rolling. After rough rolling, it enters the finishing mill and uses a multi-stand continuous rolling process to obtain strip steel, controlling the final rolling temperature to 800-880℃. 4) Laminar flow cooling After exiting the rolling mill, the strip undergoes laminar flow cooling, employing a segmented cooling method. The first stage uses laminar flow cooling with water, controlling the cooling rate to ≥100℃ / s. Then, the strip enters the air-cooling section, with an initial temperature of T±20℃, where T=700-n(α) / 0.15×50℃, and an air-cooling time of t±1s, where t=[n(α) / 0.15+p Ti [0.2]×10s; After the air cooling section ends, the second water cooling section begins, where the strip is cooled to below 200℃ at a cooling rate of ≥60℃ / s before being coiled.

Images