A method for manufacturing a low carbon equivalent high strength steel for a ship plate by TMCP

Abstract

The application discloses a TMCP preparation method of a low-carbon equivalent high-strength steel of a ship plate, and belongs to the technical field of metal material processing. The steel material comprises C, Si, Mn, Ni, Cr, Mo, Al, Nb, P, S, Fe and impurities. The preparation method comprises the following steps: smelting and continuous casting, slab heating, controlled rolling, quenching treatment and reverse phase transformation annealing. The application adopts low-carbon medium-nickel collaborative design, replaces interstitial atom carbon with substitution atom nickel as an austenite stabilizing element, introduces high-density crystal defects through controlled rolling, obtains full martensite structure in combination with offline quenching, and then is subjected to two-phase zone reverse phase transformation annealing to obtain the final structure of a tempered martensite matrix with 10% to 20% thin film-shaped reverse transformation austenite distributed thereon. The obtained steel material has high yield strength and high impact energy at-80 DEG C, and welding does not need preheating, and has high strength, excellent low-temperature toughness and excellent welding performance, and is suitable for high-value-added application scenarios such as polar ships.

Description

Technical Field

[0001] This invention relates to the field of metal material processing technology, and in particular to a method for preparing TMCP (Transformed Molecular Carbon) of low-carbon equivalent high-strength steel for ship plate steel. Background Technology

[0002] Ship plate steel is a key material for ship and marine engineering structures, requiring high strength, excellent low-temperature toughness, and good weldability. With the increasing size of ships and the growing demand for polar navigation and deep-sea operations, the performance requirements for ship plate steel are becoming increasingly stringent.

[0003] Traditional methods for preparing ship plate steel mainly combine low-carbon microalloying composition design with TMCP (Transformer Metallization Processing) to obtain a fine ferrite + bainite microstructure through controlled rolling and cooling. However, this type of method has the following technical drawbacks: Carbon segregation is difficult to eliminate: Traditional processes rely on interstitial diffusion of carbon to achieve compositional homogenization, but carbon tends to accumulate between dendrites during solidification, forming banded structures that affect low-temperature toughness and weldability.

[0004] Softening of the heat-affected zone during welding: When traditional low-carbon equivalent steel is welded at high heat input, the heat-affected zone exhibits coarsening and softening of the microstructure, leading to a decline in joint performance.

[0005] Low-temperature toughness bottleneck: Carbon-stabilized austenite is prone to instability and transformation into brittle martensite below -60℃, which makes it difficult to meet the toughness requirements of polar vessels for temperatures of -80℃ or even lower.

[0006] Limited process efficiency: In the traditional TMCP process, carbon diffusion relies on high temperature and long-term heat preservation, which limits production efficiency.

[0007] Therefore, developing a new method for preparing ship plate steel that can overcome the above-mentioned technical bottlenecks has significant engineering value and market prospects. Summary of the Invention

[0008] The purpose of this invention is to address the shortcomings of existing technologies by proposing a method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel.

[0009] To achieve the above objectives, the present invention adopts the following technical solution: This invention first proposes a low-carbon equivalent high-strength steel for ship plates, which is composed of the following chemical components by weight percentage: C≤0.02%, Si: 0.10%~0.30%, Mn: 0.50%~1.00%, Ni: 3.50%~5.00%, Cr: 0.50%~1.00%, Mo: 0.20%~0.40%, Al: 0.02%~0.05%, Nb: 0.02%~0.04%, P≤0.008%, S≤0.003%, The balance consists of Fe and unavoidable impurities.

[0010] This invention also proposes a method for preparing the aforementioned low-carbon equivalent high-strength steel TMCP for ship plate steel, comprising the following steps: S1. Smelting and continuous casting: Smelting is carried out according to the chemical composition described in claim 1, and after ladle refining and vacuum degassing, it is continuously cast into slabs, with electromagnetic stirring applied during the continuous casting process. Using Al as the final deoxidizer and Si as the auxiliary deoxidizer, the following deoxidation reactions occur: 2[Al] + 3[O] = Al2O3, [Si] + 2[O] = SiO2. The deoxidation products are removed by flotation in the LF (ladle refining furnace) refining slag (high basicity CaO-SiO2-Al2O3 slag system), ultimately controlling the inclusions of types A / B / C / D to be ≤1.0. Strict control of the Al content (0.02-0.05%) ensures both effective deoxidation and avoids the formation of large-sized brittle Al2O3 inclusions. Simultaneously, the generated nano-sized AlN can pin austenite grain boundaries and refine the grains during subsequent heating.

[0011] Phosphorus (P) tends to segregate at austenite grain boundaries, leading to low-temperature tempering brittleness. S (S) tends to form long, thin MnS inclusions, which become the nuclei for crack initiation. Both significantly degrade low-temperature toughness and must be strictly controlled. Under LF refining white slag operation, [S] reacts with CaO to form CaS, which enters the slag phase. [P] reacts with CaO in an oxidizing atmosphere to form stable calcium phosphate, which also enters the slag phase. Ultimately, P ≤ 0.008% and S ≤ 0.003% are controlled.

[0012] Ni significantly increases the solubility of hydrogen in steel. Although low-carbon design reduces the risk of hydrogen embrittlement, strict hydrogen control is still required in the preparation of thick plates to avoid defects such as hydrogen-induced white spots and delayed cracking. During the VD (vacuum argon degassing method) (vacuum degree ≤67Pa, pressure holding ≥15min), the [H] and [N] dissolved in the molten steel escape from the molten steel under vacuum, and the final control of H ≤1.5ppm is achieved.

[0013] The equilibrium distribution coefficient of substitutional alloying elements such as Ni, Cr, and Mo in γ-Fe is k0 < 1 (k0 ≈ 0.85 for Ni). During solidification, solute redistribution occurs at the solid-liquid interface. The content of alloying elements in the solid phase is low, while they continue to accumulate in the liquid phase, eventually forming segregation in the center of the billet. In particular, Ni, as a substitutional atom, has a diffusion rate in Fe that is only one ten-thousandth that of interstitial atoms, and its segregation tendency is much higher than that of ordinary carbon steel.

[0014] During continuous casting, electromagnetic stirring at 600-800A and 5-8Hz is applied to drive forced convection of the molten steel through Lorentz force. On the one hand, this breaks down the dendrites at the solid-liquid interface, refines the equiaxed crystals, and expands the equiaxed crystal region. On the other hand, it forces homogenization of solute elements such as Ni in the molten steel, suppresses interdendritic and central segregation, and achieves uniform distribution of substitution atoms across the entire cross section of the billet, laying the compositional foundation for the uniform occurrence of subsequent reverse phase transformation.

[0015] The superheat of continuous casting is controlled at 15-25℃ to further suppress excessive growth of columnar crystals, reduce dendrite segregation, and avoid fluctuations in the content of reverse-transformed austenite due to uneven composition during subsequent heat treatment.

[0016] S2. Slab heating: Heat the slab to 1180℃~1220℃ and hold for 30~50 minutes to ensure uniform heating. The slab is heated to 1180-1220℃. At this temperature, the diffusion coefficient of substitutional atoms such as Ni, Cr, and Mo in γ-Fe is significantly improved. The compositional segregation between the as-cast dendrites is eliminated through atomic diffusion, and the alloying elements are completely dissolved and uniformly distributed in the austenitic matrix, completing the basic step of "substitutional atom partitioning".

[0017] The room-temperature as-cast microstructure (mainly ferrite with a very small amount of pearlite) undergoes an α→γ phase transformation upon heating to above Ac3, completely transforming into single-phase austenite. The heating rate is controlled at 8-12℃ / min to avoid thermal stress cracking caused by excessive temperature differences between the surface and interior of thick plates; the soaking time is controlled at 30-50min to ensure sufficient solidification of alloying elements while avoiding excessive coarsening of austenite grains, with a target initial austenite grain size ≥4.

[0018] At this temperature, the carbonitride Nb(C,N) of the microalloying element Nb is completely dissolved in austenite. The dissolved Nb atoms can inhibit austenite recrystallization during subsequent rolling, thereby achieving grain refinement.

[0019] Ni is an element with weak oxidation resistance. At high temperatures, Fe is preferentially oxidized, forming a Ni-rich layer at the oxide scale / substrate interface. Severe oxidation can lead to Ni desorption from the surface, degrading the surface structure and properties. A weak oxidizing atmosphere can balance the contradiction between dephosphorization and surface oxidation. In this case, a weak oxidizing atmosphere is used, and the oxide scale thickness is controlled to be ≤1mm.

[0020] S3. Controlled rolling: This includes two stages: roughing and finishing. The initial rolling temperature for roughing is 1080℃~1120℃, and the final rolling temperature is 1000℃~1050℃, with a cumulative reduction rate of 50%~70%. The initial rolling temperature for finishing is 880℃~920℃, and the final rolling temperature is 820℃~860℃, with a cumulative reduction rate of 60%~80%. After rolling, the rolling process is followed by air cooling to room temperature. Each pass of large-scale reduction deformation elongates the austenite grains, generating numerous dislocations within the grains. The energy stored in the deformation drives dynamic recrystallization—new, distortion-free austenite nuclei nucleate and grow at the original grain boundaries and dislocation entanglements, replacing the coarse austenite grains in the as-cast state. Static recrystallization occurs between passes, further refining the grains. A cumulative reduction rate of 50-70% provides sufficient energy for deformation storage, and repeated recrystallization refines the coarse austenite grains of hundreds of micrometers in the as-cast state to tens of micrometers, achieving ultra-fine grain refinement in the first pass. Large-scale reduction deformation breaks down defects such as porosity and shrinkage cavities in the as-cast state, achieving full densification of the billet's microstructure and improving the uniformity of the core properties of thick plates.

[0021] During the finishing rolling stage, the Nb and Mo atoms dissolved in the steel significantly increase the activation energy for austenite recrystallization, dragging grain boundary migration and inhibiting austenite recrystallization after deformation. Large reduction deformation causes austenite grains to elongate significantly along the rolling direction, forming flattened austenite grains and multiplying the grain boundary area. Simultaneously, it introduces high-density dislocations into the grains, forming defect configurations such as dislocation walls, dislocation cells, and deformation bands. Ni, as a substitutional atom, has an extremely low bulk diffusion rate in the iron matrix, while its diffusion rate along dislocation channels is more than 1000 times that of bulk diffusion. The high-density dislocations introduced by finishing rolling are the core channels for the rapid diffusion of Ni atoms during subsequent reverse transformation annealing. This is crucial for achieving precise Ni distribution and uniform nucleation of reverse transformation austenite, which is the core reason why this invention requires a cumulative reduction rate of ≥70% during finishing rolling. Ni is the core element for reducing the stacking fault energy of austenite. By precisely controlling the intrinsic stacking fault energy of austenite through Ni content, austenite with low stacking fault energy is more likely to generate stacking faults and deformation twins during deformation, further increasing the defect density and nucleation sites, refining the subsequent martensitic lath structure, and achieving synergistic control between "stacking fault energy engineering" and rolling process.

[0022] This step requires air cooling to room temperature after rolling, without using online accelerated cooling. The core reason is that online rapid cooling would cause premature martensitic transformation and trigger self-tempering, leading to dislocation recovery annihilation and loss of Ni atom diffusion channels; at the same time, online cooling easily leads to uneven microstructure between the surface and interior of thick plates, affecting the uniformity of subsequent reaustitization.

[0023] S4. Quenching treatment: Heat the rolled steel plate to 880℃~920℃, hold for 20~40min, and then water quench to cool to the final cooling temperature ≤50℃ to obtain a full martensitic structure. The rolled steel plate is heated to 880-920℃ (Ac 03 The temperature is 30-50℃ (above), and the plate is held for 20-40 minutes. After rolling and air cooling, the microstructure undergoes an α→γ phase transformation, completely transforming into single-phase austenite. Alloying elements such as Ni, Cr, and Mo are fully dissolved, achieving uniformity in composition and grain size across the entire cross-section of the thick plate. The holding time is adjusted according to the plate thickness to ensure complete austenitization of the core of the thick plate and to avoid undissolved second phases affecting the uniformity of subsequent martensitic transformation.

[0024] Water quenching is performed using a roller press quenching machine, with the core cooling rate controlled at ≥20℃ / s and the final cooling temperature ≤50℃. When the austenite is cooled to below the Ms point (the Ms point of this composition system is about 350-400℃), a diffusionless shear phase transformation occurs, and the face-centered cubic γ-Fe is transformed into body-centered tetragonal α'-Fe martensite.

[0025] The cooling rate must be ≥20℃ / s, exceeding the critical quenching rate of the steel, to avoid the precipitation of non-martensitic structures such as proeutectoid ferrite and bainite; the final cooling rate must be ≤50℃ to ensure complete transformation of austenite into martensite, with a final martensite volume fraction ≥95%, corresponding to a hardness ≥450HV. Non-martensitic structures can lead to uneven nucleation during subsequent reverse deformation, forming coarse austenite and destroying the film-like morphology; therefore, the full martensitic transformation must be strictly controlled. Because the C content is ≤0.02%, the martensite has extremely low tetragonality, exhibiting typical low-carbon lath martensite rather than the twinned martensite of high-carbon steel. Multiple martensite lath bundles are formed within an austenite grain, with parallel nanoscale martensite laths within each bundle, large-angle grain boundaries between the laths, and high-density dislocations (102) preserved within the laths. 15 -10 16 m -2 This provides nucleation sites for the subsequent reverse phase transition and channels for Ni atom diffusion.

[0026] S5. Reverse phase transformation annealing: The quenched steel plate is heated to 620℃~660℃ and held for 30~90min. Then it is air-cooled or slowly cooled to room temperature at a cooling rate of ≤10℃ / min to obtain the final microstructure with a thin film of reverse austenite distributed on the tempered martensite matrix.

[0027] The heating rate is controlled to be ≤5℃ / min. When the temperature rises from room temperature to the target annealing temperature, the core change is the recovery of martensite: when heated to above 400℃, the high-density dislocations in the quenched martensite recover, and the dislocation tangles decompose into regular dislocation cells. The dislocation density decreases moderately but retains the core diffusion channels. The martensite lath boundary structure becomes clearer, providing a preferred site for the nucleation of austenite in reverse transformation. Slow heating also avoids deformation and internal stress caused by excessive temperature difference between the surface and the interior of the thick plate.

[0028] During the heat preservation process, two major processes occur simultaneously: martensitic tempering and the nucleation, growth, and compositional stabilization of reverse-transformed austenite. The reverse transformation is a diffusion-type transformation, which is completely different from the diffusionless martensitic transformation during quenching.

[0029] Reversed austenite preferentially nucleates at martensite lath boundaries, lath bundle boundaries, original austenite grain boundaries, and dislocation walls. The mechanism is that these locations have high defect density, large distortion energy, and low nucleation activation energy; at the same time, Ni, as an austenite stabilizing element, spontaneously diffuses towards the nucleation sites, lowering the free energy barrier for γ-phase nucleation.

[0030] The quenched lath martensite has a large number of two-dimensional lath boundaries. The austenite nucleates uniformly on the two-dimensional interface, rather than at the three-dimensional grain boundary junction. Structurally, this determines that the growth morphology of austenite is a thin film extending along the lath boundaries, rather than an equiaxed shape, which lays the structural foundation for obtaining nanoscale austenite with a thickness ≤100nm.

[0031] After austenite nucleation, it grows in two dimensions along the martensite lath boundaries. At the same time, the core Ni atoms are redistributed: the solubility of Ni in the face-centered cubic γ phase is much higher than that in the body-centered cubic α' phase. Therefore, Ni atoms diffuse rapidly through dislocation channels and are continuously enriched from the surrounding martensite matrix into the reverse-transformed austenite, so that the Ni content in the austenite can reach more than 10%, which is much higher than the average Ni content of 3.5-5% in the matrix.

[0032] The matching mechanism between temperature and Ni content: The higher the Ni content, the lower the γ-phase free energy, the greater the driving force for the reverse phase transformation, and the higher the required temperature. Therefore, 3.5-4.2% Ni corresponds to annealing at 620-640℃, and 4.2-5.0% Ni corresponds to annealing at 640-660℃. At lower temperatures, Ni diffusion is insufficient, the austenite nucleation rate is low, the final content is <5%, and the improvement in toughness is limited; at higher temperatures, Ni diffusion is too rapid, the austenite breaks through the thin-film morphology and coarsens into equiaxed crystals (>5μm), Ni distribution is uneven, and stability decreases significantly.

[0033] Mechanism of heat preservation time control: A heat preservation time of 30-90 minutes provides sufficient time for Ni atoms to diffuse fully and complete the compositional stabilization of austenite. If the time is too short, Ni distribution is insufficient, the Ms point of austenite is higher than room temperature, and it will transform into martensite upon cooling and cannot be retained; if the time is too long, the austenite coarsens, stability decreases, and the martensitic matrix softens excessively, resulting in strength loss.

[0034] Simultaneous martensitic tempering occurs during the reverse phase transformation. Because the C content is ≤0.02%, no brittle carbides such as cementite are precipitated during the tempering process. Only dislocations are further recovered, internal stress is eliminated, and the lath structure is stabilized, forming a pure tempered lath martensitic matrix. This ensures both matrix strength and extremely high intrinsic low-temperature toughness, avoiding the low-temperature embrittlement caused by carbide precipitation in traditional quenched and tempered steel.

[0035] After the heat preservation period, the material is slowly cooled to room temperature at a rate of ≤10℃ / min. The core objective is to ensure the complete preservation of the stabilized reversible austenite at room temperature. Through Ni enrichment during the heat preservation stage, the Ni content in the reversible austenite is ≥10%, and the Ms point has dropped below room temperature, preventing martensitic transformation during cooling. Simultaneously, the size effect of the nanoscale thin-film austenite further lowers the Ms point, enhancing thermodynamic stability and ultimately achieving stable retention of 10-20% austenite. Excessive cooling can cause thermal stress leading to steel plate deformation and easily trigger strain-induced martensitic transformation, resulting in a loss of austenite content. Slow cooling further eliminates internal stress, ensuring the flatness of the steel plate while allowing the austenite to complete its final compositional stabilization. A protective atmosphere of N2 + 5-10% H2 with a dew point ≤-40℃ is used to avoid Ni desorption caused by preferential Fe oxidation on the steel plate surface at high temperatures, ensuring the uniformity of microstructure and properties between the steel plate surface and core.

[0036] Preferably, in step S1, the current intensity of the electromagnetic stirring is 600A to 800A, and the frequency is 5Hz to 8Hz.

[0037] Preferably, in step S2, the heating rate of the slab is controlled at 8℃ / min to 12℃ / min.

[0038] Preferably, in step S3, the cumulative reduction rate in the finishing rolling stage is ≥70%, and the finishing rolling temperature is controlled within the range of 820℃~860℃.

[0039] Preferably, in step S4, water quenching is performed using a roller press quenching machine with a cooling rate ≥20℃ / s and a temperature difference in the thickness direction of the steel plate ≤50℃.

[0040] Preferably, in step S5, the reverse phase transformation annealing is carried out in a roller hearth continuous heat treatment furnace with a protective atmosphere, which is a mixture of nitrogen and hydrogen, wherein the hydrogen component is 5% to 10% and the furnace dew point is ≤-40℃.

[0041] Preferably, in step S5, the reverse phase transformation annealing temperature is precisely controlled according to the nickel content: when the Ni content is 3.50% to 4.20%, the annealing temperature is 620℃ to 640℃; when the Ni content is 4.20% to 5.00%, the annealing temperature is 640℃ to 660℃.

[0042] Preferably, in step S5, the reverse-transformed austenite obtained after reverse phase transformation annealing is in the form of a thin film, distributed at the boundaries of the martensite laths, and the film thickness is ≤100nm.

[0043] Preferably, the steel has a yield strength ≥460MPa, an average impact energy of -80℃ ≥200J, and does not require preheating treatment during welding.

[0044] Compared with the prior art, the beneficial effects of the present invention are: 1. A synergistic design of low carbon (C≤0.02%) and medium nickel (3.5%~5.0%) is adopted, with substitutional nickel replacing interstitial carbon as an austenite stabilizing element, eliminating the risk of carbide precipitation and hydrogen embrittlement, while significantly improving the low-temperature toughness limit of the steel.

[0045] 2. By using the process path of "controlled rolling + offline quenching + reverse transformation annealing", the high-density crystal defects introduced during the rolling stage are used to accelerate the diffusion of nickel during the reverse transformation. 10% to 20% of thin film reverse transformation austenite can be obtained by holding at 620℃ to 660℃ for 30 to 90 minutes, thus achieving a synergy of high strength and excellent low-temperature toughness.

[0046] 3. Since the steel does not contain carbides and the reverse-transformed austenite has high thermal stability, the weld heat-affected zone has no tendency to become embrittled, enabling welding processes that do not require preheating, supporting high heat input welding, and significantly improving shipbuilding efficiency.

[0047] 4. The steel maintains excellent impact toughness within a temperature range of -80℃ to -196℃, making it suitable for high-value-added applications such as polar vessels and liquefied natural gas carriers. Attached Figure Description

[0048] Figure 1 This is a process flow diagram of the low-carbon equivalent high-strength steel proposed in this invention. Detailed Implementation

[0049] The technical solutions in the embodiments of the present invention will be clearly and completely described below in conjunction with existing known technologies. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments.

[0050] Example 1: A low-carbon equivalent high-strength steel for ship plates, composed of the following chemical components by weight percentage: C≤0.02%, Si: 0.10%, Mn: 1.00%, Ni: 3.50%, Cr: 1.00% Mo: 0.20%, Al: 0.05%, Nb: 0.02%, P≤0.008%, S≤0.003%, The balance consists of Fe and unavoidable impurities.

[0051] TMCP preparation method for low-carbon equivalent high-strength steel of ship plate, such as Figure 1 As shown, it includes the following steps: S1. Smelting and continuous casting: Smelting is carried out according to the chemical composition described in claim 1, and after ladle refining and vacuum degassing, it is continuously cast into slabs, with electromagnetic stirring applied during the continuous casting process. S2. Slab heating: Heat the slab to 1220℃ and hold for 30 minutes to ensure uniform heating. S3. Controlled rolling: This includes two stages: roughing and finishing. The roughing temperature is 1120℃ at the start and 1050℃ at the finish, with a cumulative reduction of 67%. The finishing temperature is 880℃ at the start and 820℃ at the finish, with a cumulative reduction of 82%. After rolling, the rolling process is followed by air cooling to room temperature. S4. Quenching treatment: Heat the rolled steel plate to 920℃ and hold for 20 minutes, then quench it with water until the final cooling temperature is ≤50℃ to obtain a full martensitic structure. S5. Reverse phase transformation annealing: The quenched steel plate is heated to 620℃ and held for 90 min. Then it is air-cooled or slowly cooled to room temperature at a cooling rate of ≤10℃ / min to obtain the final microstructure with a thin film of reverse austenite distributed on the tempered martensite matrix.

[0052] In S1, the electromagnetic stirring current intensity is 600A and the frequency is 8Hz.

[0053] In step S2, the heating rate of the slab is controlled at 8°C / min.

[0054] In S3, the cumulative reduction rate in the finishing rolling stage is ≥70%, and the finishing rolling temperature is controlled within the range of 820℃~860℃.

[0055] In S4, water quenching is performed using a roller press quenching machine with a cooling rate ≥20℃ / s and a temperature difference in the thickness direction of the steel plate ≤50℃.

[0056] In step S5, the reverse phase transformation annealing is carried out in a roller hearth continuous heat treatment furnace with a protective atmosphere, which is a mixture of nitrogen and hydrogen, wherein the hydrogen component is 10% and the furnace dew point is ≤-40℃.

[0057] In S5, the reverse-transformed austenite obtained after reverse phase transformation annealing is in the form of a thin film, distributed at the boundaries of martensite laths, with a film thickness ≤100nm.

[0058] Example 2: The preparation method is the same as in Example 1, except that, by weight percentage, it consists of the following chemical components: C≤0.02%, Si: 0.20%, Mn: 0.75%, Ni: 4.20%, Cr: 0.75% Mo: 0.30%, Al: 0.035%, Nb: 0.03%, P≤0.008%, S≤0.003%, The balance consists of Fe and unavoidable impurities.

[0059] Includes the following steps: S1. Smelting and continuous casting: Smelting is carried out according to the chemical composition described in claim 1, and after ladle refining and vacuum degassing, it is continuously cast into slabs, with electromagnetic stirring applied during the continuous casting process. S2. Slab heating: Heat the slab to 1200℃ and hold for 40 minutes to ensure uniform heating. S3. Controlled rolling: This includes two stages: roughing and finishing. The roughing temperature is 1100℃ at the start and 1025℃ at the finish, with a cumulative reduction of 61%. The finishing temperature is 900℃ at the start and 840℃ at the finish, with a cumulative reduction of 75%. After rolling, the rolling process is followed by air cooling to room temperature. S4. Quenching treatment: Heat the rolled steel plate to 900℃, hold for 30 minutes, and then water quench to a final cooling temperature of ≤50℃ to obtain a fully martensitic structure. S5. Reverse phase transformation annealing: The quenched steel plate is heated to 640℃ and held for 60 min. Then it is air-cooled or slowly cooled to room temperature at a cooling rate of ≤10℃ / min to obtain the final microstructure with a thin film of reverse austenite distributed on the tempered martensite matrix, with a volume fraction of 10% to 20%.

[0060] In S1, the electromagnetic stirring current intensity is 700A and the frequency is 6Hz.

[0061] In step S2, the heating rate of the slab is controlled at 10℃ / min.

[0062] Example 3: The preparation method is the same as in Example 1, except that, by weight percentage, it consists of the following chemical components: C≤0.02%, Si: 0.30%, Mn: 0.50%, Ni: 5.00% Cr: 0.50% Mo: 0.40%, Al: 0.02%, Nb: 0.04%, P≤0.008%, S≤0.003%, The balance consists of Fe and unavoidable impurities.

[0063] Includes the following steps: S1. Smelting and continuous casting: Smelting is carried out according to the chemical composition described in claim 1, and after ladle refining and vacuum degassing, it is continuously cast into slabs, with electromagnetic stirring applied during the continuous casting process. S2. Slab heating: Heat the slab to 1180℃ and hold for 50 minutes to ensure uniform heating. S3. Controlled rolling: This includes two stages: roughing and finishing. The roughing temperature is 1080℃ at the start and 1000℃ at the finish, with a cumulative reduction of 56%. The finishing temperature is 920℃ at the start and 860℃ at the finish, with a cumulative reduction of 72%. After rolling, the rolling process is followed by air cooling to room temperature. S4. Quenching treatment: Heat the rolled steel plate to 880℃ and hold for 40 minutes, then quench it with water until the final cooling temperature is ≤50℃ to obtain a full martensitic structure. S5. Reverse phase transformation annealing: The quenched steel plate is heated to 660℃ and held for 30 min. Then it is air-cooled or slowly cooled to room temperature at a cooling rate of ≤10℃ / min to obtain the final microstructure with a thin film of reverse austenite distributed on the tempered martensite matrix.

[0064] In S1, the electromagnetic stirring current intensity is 800A and the frequency is 5Hz.

[0065] In step S2, the heating rate of the slab is controlled at 12℃ / min.

[0066] The following comparison model was also set: Comparative Example 1: Based on Example 2, the difference is that: C: 0.07%, Si: 0.25%, Mn: 1.20%, Ni: 3.50%, Cr: 0.20%, Mo: 0.10%, Al: 0.030%, P: 0.012%, S: 0.008%, with the balance being Fe and unavoidable impurities, and the rest being the same as Example 2.

[0067] Comparative Example 2: Based on Example 2, the difference is that only the C content was adjusted to 0.06%, and the rest is the same as Example 2.

[0068] Comparative Example 3: Based on Example 2, the only difference is that the Ni content is adjusted to 2.0%, and the rest is the same as Example 2.

[0069] Comparative Example 4: Based on Example 2, the difference is that only the S4 full martensitic quenching process is omitted, and after S3 rolling and air cooling to room temperature, S5 reverse phase transformation annealing is directly carried out. The rest is the same as Example 2.

[0070] Comparative Example 5: Based on Example 2, the difference is that only the reverse phase transformation annealing temperature of S5 is adjusted to 680°C, and the rest is the same as Example 2.

[0071] Comparative Example 6: Based on Example 2, the difference is that only the cumulative reduction rate of S3 finishing rolling is adjusted to 50%, and the rest is the same as Example 2.

[0072] Performance Testing: A 20mm thick industrial ship plate was used as the test object. The room temperature tensile properties of the invention were tested according to GB_T228_1-2021 Metallic Materials - Tensile Testing - Part 1 - Room Temperature Test Method; the low temperature impact toughness of the invention was tested according to DNV-OS-B101-2025 Offshore Metallic Materials Standard; and the hardness and thickness-direction uniformity of the invention were tested according to GB / T 4340.1-2024 Metallic Materials - Vickers Hardness Test - Part 1: Test Method. The results are shown below: Table 1. Test results of various properties of low carbon equivalent high strength ship plate steel Data Analysis: All three sets of embodiments fully achieved the design objectives of this invention, with yield strengths of 470–510 MPa, tensile strengths of 590–640 MPa, and average V-notch impact energy of 210–245 J at -80℃, comprehensively meeting the performance requirements of high-strength, low-temperature ship plates. Embodiment 2 is the optimal solution, achieving a uniform, thin-film reverse-transformed austenite of 16 vol% and an average thickness of 60 nm through precise matching of composition and process. The average impact energy at -80℃ reaches 245 J, a 166% improvement over the industry-standard 3.5Ni steel. Simultaneously, the cold cracking sensitivity coefficient Pcm is only 0.18%, enabling preheating-free welding, with a thickness-direction hardness difference of ≤5 HV, and excellent microstructure uniformity.

[0073] The comparative data accurately verified the irreplaceable nature of the core technical features of this invention. In terms of composition, Comparative Example 2 had a carbon content exceeding the standard by 0.06%, leading to brittle carbide precipitation and a 72% drop in impact energy at -80℃; Comparative Example 3 had a nickel content reduced to 2.0%, with only 4 vol% of stable reverse austenite, resulting in complete loss of the TRIP toughening effect and an impact energy only 17% of Example 2. In terms of process, Comparative Example 4 omitted full martensitic quenching, Comparative Example 5 had an excessive annealing temperature, and Comparative Example 6 had insufficient finishing rolling reduction, all of which failed to obtain continuous thin-film reverse austenite, resulting in a drop in low-temperature impact energy exceeding 65%, along with problems such as deterioration of microstructure uniformity and a sharp decrease in matrix strength.

[0074] The above description is only a preferred embodiment of the present invention, but the scope of protection of the present invention is not limited thereto. Any equivalent substitutions or modifications made by those skilled in the art within the scope of the technology disclosed in the present invention, based on the technical solution and inventive concept of the present invention, should be covered within the scope of protection of the present invention.

Claims

1. A low-carbon equivalent high-strength steel for ship plates, characterized in that, It consists of the following chemical components by weight percentage: C≤0.02%， Si: 0.10%~0.30%, Mn: 0.50%~1.00%, Ni: 3.50%~5.00%, Cr：0.50%～1.00%， Mo: 0.20%~0.40%, Al：0.02%～0.05%， Nb: 0.02%~0.04%, P≤0.008%， S≤0.003%， The balance consists of Fe and unavoidable impurities.

2. A method for preparing low-carbon equivalent high-strength steel TMCP as described in claim 1, characterized in that, Includes the following steps: S1. Smelting and continuous casting: Smelting is carried out according to the chemical composition described in claim 1, and after ladle refining and vacuum degassing, it is continuously cast into slabs, with electromagnetic stirring applied during the continuous casting process. S2. Slab heating: Heat the slab to 1180℃~1220℃ and hold for 30~50 minutes to ensure uniform heating. S3. Controlled rolling: This includes two stages: roughing and finishing. The initial rolling temperature for roughing is 1080℃~1120℃, and the final rolling temperature is 1000℃~1050℃, with a cumulative reduction rate of 50%~70%. The initial rolling temperature for finishing is 880℃~920℃, and the final rolling temperature is 820℃~860℃, with a cumulative reduction rate of 60%~80%. After rolling, the rolling process is followed by air cooling to room temperature. S4. Quenching treatment: Heat the rolled steel plate to 880℃~920℃, hold for 20~40min, and then water quench to cool to the final cooling temperature ≤50℃ to obtain a full martensitic structure. S5. Reverse phase transformation annealing: The quenched steel plate is heated to 620℃~660℃ and held for 30~90min. Then it is air-cooled or slowly cooled to room temperature at a cooling rate of ≤10℃ / min to obtain the final microstructure with a thin film of reverse austenite distributed on the tempered martensite matrix.

3. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In S1, the current intensity of the electromagnetic stirring is 600A to 800A, and the frequency is 5Hz to 8Hz.

4. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In step S2, the heating rate of the slab is controlled at 8℃ / min to 12℃ / min.

5. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In S3, the cumulative reduction rate in the finishing rolling stage is ≥70%, and the finishing rolling temperature is controlled within the range of 820℃~860℃.

6. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In S4, water quenching is performed using a roller press quenching machine with a cooling rate ≥20℃ / s and a temperature difference in the thickness direction of the steel plate ≤50℃.

7. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In step S5, the reverse phase transformation annealing is carried out in a roller hearth continuous heat treatment furnace with a protective atmosphere, which is a mixture of nitrogen and hydrogen, wherein the hydrogen component is 5% to 10% and the furnace dew point is ≤-40℃.

8. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In step S5, the reverse phase transformation annealing temperature is precisely controlled according to the nickel content: when the Ni content is 3.50% to 4.20%, the annealing temperature is 620℃ to 640℃; when the Ni content is 4.20% to 5.00%, the annealing temperature is 640℃ to 660℃.

9. The method for preparing low-carbon equivalent high-strength steel TMCP for ship plate steel according to claim 2, characterized in that, In S5, the reverse-transformed austenite obtained after reverse phase transformation annealing is in the form of a thin film, distributed at the boundaries of martensite laths, with a film thickness ≤100nm.

10. The low-carbon equivalent high-strength steel for ship plates according to claim 1, characterized in that, The steel has a yield strength ≥460MPa, an average impact energy of ≥200J at -80℃, and requires no preheating treatment during welding.

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