Steel plate with strong anti-collision capability, high toughness, high ductility and super-large heat input and manufacturing method thereof

Abstract

The invention discloses a high-toughness and high-ductility steel plate capable of achieving super-large heat input and a manufacturing method thereof. The steel plate comprises, by weight, 0.09%-0.15% of C, smaller than or equal to 0.20% of Si, 1.30%-1.60% of Mn, smaller than or equal to 0.013% of P, smaller than or equal to 0.0030% of S, 0.008%-0.022% of Nb, 0.008%-0.016% of Ti, 0.005%-0.015% of Als, 0.0055%-0.0085% of N, 0.0012%-0.0030% of B, 0.0010%-0.0035% of Ca and the balance Fe and inevitable impurities. According to the invention, precious elements such as Ni, Cu and Mo are not added, an alloy system treated by medium C, ultra-low (Si + Al), high Mn, high N and micro (Ti + B) is adopted as a basis, the solidification co-segregation index is controlled to be less than or equal to 2.52, Ti / N is less than or equal to 2.0, Ca treatment is carried out, the Ca / S ratio is 1.0-3.0, and (% Ca) * (% S) 0.28 is equal to 1.0 * 10 <-3 >-3.0 * 10 <-3 >. The TMCP technology is optimized, high strength, high impact characteristic, high toughness (especially high crack arrest characteristic) and excellent weldability are obtained, and ultra-large heat input welding can be achieved.

Description

Technical Field

[0001] The present invention relates to high-strength steel, and in particular to a steel plate with high toughness, high ductility, strong impact resistance and ultra-large heat input and a manufacturing method thereof. Background Art

[0002] As we all know, low carbon (high strength) low alloy steel is one of the most important engineering structural materials, widely used in oil and gas pipelines, offshore platforms, shipbuilding, bridge structures, boiler containers, building structures, automobile industry, railway transportation and machinery manufacturing; the performance of low carbon (high strength) low alloy steel depends on its chemical composition and manufacturing process, among which strength, toughness, plasticity and weldability are the most important properties of low carbon (high strength) low alloy steel, which is ultimately determined by the microstructure state of the finished steel.

[0003] With the continuous development of metallurgical technology and on-site control technology, people have put forward higher requirements on the toughness, plasticity and weldability of high-strength steel; that is, the steel plate has high strength, high elongation (especially high uniform elongation), high crack arrest characteristics (i.e. resistance to brittle fracture and plastic instability fracture) at low temperature, while the steel plate has excellent welding performance and can be welded with high heat input and high efficiency; and under the condition of lower manufacturing cost, the comprehensive mechanical properties and performance of the steel plate are greatly improved to reduce the amount of steel and save costs, reduce the weight, stability and safety of the steel structure itself, and more importantly, to further improve the cold and hot workability of high-strength steel and the safety and reliability during service.

[0004] At present, Japan, South Korea, the European Union and North America have set off a research climax on the development of a new generation of high-performance steel materials, trying to combine the alloy combination design optimization with the new generation of online deformation heat treatment process technology to obtain better complex phase microstructure matching, achieve ultra-fine microstructure, and evenly distributed CELL structure dislocation configuration. Without adding precious alloy elements (such as Cu, Ni, Mo, etc.), through the optimization of component combination design and matching with the new generation of TMCP process technology, better organizational morphology, size and substructure are obtained, thereby obtaining higher toughness / plasticity, higher elongation and better weldability.

[0005] In the prior art, when manufacturing thick steel plates with a yield strength of ≥355 MPa and a low-temperature impact toughness of ≥34 J at -60°C, a certain amount of Ni or (Cu+Ni) elements (≥0.30%) are generally added to the steel (The Firth (1986) international Symposium and Exhibit on Offshore Mechanics and Arctic Engineering, 1986, Tokyo, Japan, 354; "DEVELOPMENTS IN MATERIALS FOR ARCTIC OFFSHORE STRUCTURES"; "Structural Steel Plates for Arctic Use Produced by Multipurpose Accelerated Cooling System" (Japanese), "Kawasaki Steel Technology Report", 1985, No. 1 68-72; "Application of Accelerated Cooling For Producing 360 MPa Yield Strength Steel Plates of up to 150 mm in Thickness with Low Carbon Equivalent", Accelerated Cooling RolledSteel, 1986, 209~219; "High Strength Steel Plates For Ice-Breaking VesselsProduced by Thermo-Mechanical Control Process", Accelerated Cooling RolledSteel, 1986, 249~260; "420MPa Yield Strength Steel Plate with Superior FractureToughness for Arctic Offshore Structures", Kawasaki steel technical report, 1999, No. 40, 56; "420MPa and 500MPa Yield Strength Steel Plate with High HAZtoughness Produced by TMCP for Offshore Structure", Kawasaki steel technicalreport, 1993, No.29, 54; "Toughness Improvement in Bainite Structure by Thermo-Mechanical Control Process" (Japanese) "Sumitomo Metal", Vol.50, No.1 (1998), 26; "Steel Plate for Offshore Platform Structure Used in Ice Sea Areas" (Japanese), "Steel Research", 1984, No. 314, 19-43], to ensure that the base steel plate has excellent low-temperature toughness. When welding with a line energy of <100KJ / cm, the toughness of the heat-affected zone (HAZ) can also reach -60℃Akv≥34J; However, when welding with a large line energy (≥200KJ / cm), the low-temperature toughness of the heat-affected zone (HAZ) is generally difficult to achieve, and the low-temperature toughness of the heat-affected zone (HAZ) is seriously deteriorated. .

[0006] A large number of patent documents only explain how to achieve low-temperature toughness of the base steel plate, but there are few descriptions on how to obtain excellent low-temperature toughness of the heat affected zone (HAZ) under welding conditions, especially how to ensure the low-temperature toughness of the heat affected zone (HAZ) when using ultra-large input energy welding. In addition, in order to ensure the low-temperature toughness of the steel plate, a certain amount of Ni or Cu+Ni elements are generally added to the steel, and the low-temperature toughness of the heat affected zone (HAZ) of the steel plate with large input energy welding can rarely reach -60°C. (See Japanese Patents Sho 63-93845, Sho 63-79921, Sho 60-258410, Special Ping Kai 4-285119, Special Ping Kai 4-308035, Ping 3-264614, Ping 2-250917, Ping 4-143246, US Patent 4855106, US Patent 5183198, US Patent 4137104).

[0007] At present, the main technologies for improving the low temperature toughness of the heat affected zone (HAZ) of high-energy-input welded steel plates include oxide metallurgy technology (see U.S. Patent 4629505, WO 01 / 59167 A1), Ti-B treatment technology (see Japanese Patent Publication No. 59-2733, Publication No. 59-3537, Patent Application No. 56-127555, Patent Application No. 56-209177) and ultra-low N-high Al-micro-Ti treatment ("Journal of the Japanese Welding Society", 1982, Vol. 51(2), 118).

[0008] China Baosteel has developed a series of high heat input welding various low temperature steel plates (such as Chinese patents ZL201410300713.X, ZL201310244712.3, ZL201310244706.8, ZL201310124065.2, ZL201310244713.8, ZL201210209637.2, ZL201410815614.5 and ZL201710183350.X, ZL201910149978.7). In order to ensure the low-temperature toughness of the heat-affected zone of high heat input welding, a certain amount of precious alloy elements Cu and Ni are added to the steel plate (especially the ultra-thick steel plate). Although the high heat input welding performance of the steel plate is good, the toughness of the heat-affected zone of the welding, especially the low-temperature toughness of the heat-affected zone of the steel plate welding, is not very stable; especially the crack arrest properties under -40°C conditions (Kca (-40°℃) ≥8000N / mm 3 / 2 ) Not only the base steel plate and welding heat affected zone (HAZ) are not involved, not to mention the impact resistance and ultra-large heat input welding (welding heat input ≥ 300kJ / cm). Summary of the invention

[0009] The object of the present invention is to provide a steel plate with high toughness, high ductility and ultra-large heat input with strong impact resistance and a manufacturing method thereof. While obtaining excellent low-temperature toughness of the base steel plate, especially high crack arrest characteristics, the steel plate has excellent impact resistance (high uniform elongation, high fracture elongation) and weldability, and under ultra-large heat input welding conditions, the low-temperature toughness and crack arrest characteristics of the welded HAZ are also excellent; the contradiction between low alloy cost, low carbon equivalent and high strength, impact resistance, excellent low-temperature toughness (especially high crack arrest characteristics), and excellent weldability (especially ultra-large heat input weldability) is successfully solved; that is, the key factors affecting the high strength, impact resistance, excellent low-temperature toughness, high crack arrest characteristics, ultra-large heat input weldability and low-cost manufacturing of the steel plate are comprehensively considered; its yield strength is ≥355MPa, tensile strength is ≥490MPa, uniform elongation is ≥22%, elongation after fracture is ≥40% (i.e., impact resistance), -40℃ lateral impact toughness KV 2 ≥100J, low temperature crack arrest steel plate capable of high heat input welding, KV of welding HAZ 2 (-40℃)≥70J; especially suitable for offshore wind power, low-temperature pressure vessels, offshore platforms and bridge steel, etc.

[0010] To achieve the above object, the technical solution of the present invention is:

[0011] The present invention adopts a simple combination design of alloy elements, does not add Ni, Cu and Mo precious elements, adopts a medium C-ultra-low (Si+Al)-high Mn-high N-micro (Ti+B) treated alloy system as a basis, and controls the solidification co-segregation index (%C)0.5 ×[10.2(%P)+30.7(%S)+3.3(%Mn)]×[7.8(%Si)+5.6(%Als)]≤2.52, [1.27(%B)] / [(%N)-0.292(%Ti)-0.518(%Als)-0.147(%Nb)]≤61.5, Ti / N≤2.0, and Ca treatment, Ca / S ratio controlled at 1.0-3.0, and (%Ca)×(%S) 0.28 ≤1.5×10 -3 and other metallurgical techniques; optimize the TMCP (Thermo-mechanical control process) process, i.e. [ξ×(T 开冷2 )×(DI OL )×(V c2 )] / {H×(T 开轧 )×[Ar 3 -(T 开冷2 )]×(T 停冷2 )}≥7.6×10 -3 The microstructure of the finished steel plate is made of uniform and fine ferrite + a small amount of dispersed lower bainite, with an average grain size of about 20μm, achieving high strength, high impact properties, high toughness (especially high crack arrest properties) and excellent weldability and capable of ultra-large heat input welding.

[0012] Specifically, the steel plate with high toughness, high ductility and ultra-large heat input having strong impact resistance according to the present invention has the following components in percentage by weight:

[0013] C: 0.09~0.15%

[0014] Si: ≤0.20%

[0015] Mn: 1.30%~1.60%

[0016] P: ≤0.013%

[0017] S: ≤0.0030%

[0018] Nb: 0.008%~0.022%

[0019] Ti: 0.008%~0.016%

[0020] Als: 0.005%~0.015%

[0021] N: 0.0055%~0.0085%

[0022] B: 0.0012%~0.0030%

[0023] Ca: 0.0010%~0.0035%

[0024] The balance includes Fe and other unavoidable inclusions; and the contents of the above elements must simultaneously satisfy the following relationship:

[0025] Solidification co-segregation index:

[0026] (%C) 0.5 ×[10.2(%P)+30.7(%S)+3.3(%Mn)]×[7.8(%Si)+5.6(%Als)]≤2.52;

[0027] [1.27(%B)] / [(%N)-0.292(%Ti)-0.518(%Als)-0.147(%Nb)]≤61.5;

[0028] Ti / N≤2.0;

[0029] Ca treatment, Ca / S ratio = 1.0-3.0, (%Ca) × (%S) 0.28 =1.0×10 －3 ~3.0×10 －3 ;

[0030] Furthermore, the balance is Fe and other inevitable inclusions.

[0031] The microstructure of the steel plate of the present invention is uniform and fine equiaxed ferrite + a small amount of dispersed bainite.

[0032] The steel plate of the present invention has a yield strength of ≥355MPa, a tensile strength of ≥490MPa, a uniform elongation of ≥22%, a post-break elongation, i.e., an anti-impact property of ≥40%, and a transverse impact toughness of KV at -40°C. 2 ≥100J, KV of welding HAZ 2 (-40℃)≥70J.

[0033] In the composition design of the steel plate of the present invention:

[0034] C has a great influence on the strength of steel, impact resistance steel plates (high uniform elongation, high elongation at break), low-temperature toughness and weldability. From the perspective of improving the low-temperature toughness and weldability of steel, it is hoped that the C content in the steel is controlled at a low level. However, from the perspective of the strength of steel and the control of the complex phase structure during the production process, the C content should not be too low. Especially for impact resistance steel plates, too low C content (<0.090%) causes serious imbalance in the phase ratio, morphological characteristics, phase distribution and relative size of the grains of each component phase of the complex phase structure, seriously deteriorating the impact resistance and low-temperature toughness of the steel plate. In addition, it is not conducive to the strength-toughness / strength-plasticity matching of the steel plate. When the C content increases, although it is beneficial to the refinement of the microstructure of the steel plate, it damages the weldability of the steel plate, especially under ultra-large heat input welding conditions. Due to the serious coarsening of the grains in the heat affected zone (HAZ) and the slow cooling rate during the welding thermal cycle cooling process, abnormal structures such as coarse ferrite side plate strips (FSP), Widmanstatten structure (WF), and upper bainite (Bu) are easily formed in the heat affected zone (HAZ), and the number and size of MA islands increase, which seriously damages the toughness of the heat affected zone (HAZ); in addition, too high C content not only deteriorates the plasticity (uniform elongation, elongation at break) of the steel plate, which is not conducive to the impact resistance; but also deteriorates the phase ratio, morphological characteristics, phase distribution and grain size matching of the composite phase structure, which in turn further deteriorates the impact resistance and low-temperature toughness of the steel plate; therefore, the C content should not be higher than 0.15%.

[0035] Si promotes the deoxidation of molten steel and can improve the strength of steel plates. However, for molten steel deoxidized by Al, the deoxidation effect of Si is not significant. Although Si can improve the strength of steel plates, Si increases the degree of solidification segregation of molten steel, seriously damaging the low-temperature toughness and weldability of steel plates. Especially under ultra-large heat input welding conditions, Si not only promotes the formation of massive MA islands, but also the formed MA islands are large in size and unevenly distributed, seriously damaging the toughness of the welding heat affected zone (HAZ). Therefore, the Si content in steel should be controlled as low as possible. Considering the economy and operability of the steelmaking process, the Si content is controlled below 0.20%.

[0036] As the most important alloying element in steel, Mn not only improves the strength of the steel plate, but also has the functions of expanding the austenite phase and reducing Ar 3The internal Mn content in the steel should not be less than 1.30% to improve the low-temperature toughness of the steel plate by reducing the temperature, refining the uniform grain size, and promoting the formation of bainite in the TMCP process to improve the strength of the steel plate. Mn is prone to segregation during the solidification of molten steel, especially when the Mn content is too high (when the Mn content is >1.60%). It is not only easy to co-segregate with elements such as C, P, and S, but also aggravates the segregation and looseness of the center of the continuous casting billet. Severe center segregation of the continuous casting billet is prone to form abnormal tissue during the subsequent controlled rolling and welding process. In addition, excessive Mn content will also form coarse MnS particles. Such coarse MnS particles extend along the rolling direction during hot rolling, seriously deteriorating the base steel plate (especially the transverse direction) and the welding heat affected zone (HAZ) [ Especially under the condition of ultra-large heat input welding, the impact toughness (i.e., promoting the formation of coarse upper bainite, increasing the size and number of blocky MA) and the lamellar tearing resistance are poor; in addition, too high Mn content will also increase the hardenability of steel, increase the welding cold crack sensitivity coefficient (Pcm) in steel, and affect the welding processability of steel (when welding with a small heat input, it is easy to form brittle and hard structures such as martensite; when welding with ultra-large heat input, it is easy to form coarse ferrite side plate strips (FSP), Widmanstatten structure (WF), upper bainite (Bu) and other abnormal structures). Therefore, the upper limit of Mn content in steel cannot exceed 1.60%.

[0037] As a harmful inclusion in steel, P has a great detrimental effect on the mechanical properties of steel, especially low-temperature impact toughness and weldability. In theory, the lower the better, but considering the operability and cost of steelmaking, for steel plates that require ultra-large heat input welding and -40℃ toughness, the P content needs to be controlled at ≤0.013%.

[0038] As a harmful inclusion in steel, S has a great detrimental effect on the low-temperature toughness of steel. More importantly, S combines with Mn in steel to form MnS inclusions. During the hot rolling process, the plasticity of MnS causes MnS to extend along the rolling direction, forming a MnS inclusion belt along the rolling direction, which seriously damages the impact toughness (especially the transverse impact toughness), Z-direction performance and weldability of the steel plate. At the same time, S is also the main element that produces hot brittleness during hot rolling. Theoretically, the lower the requirement, the better, but considering the operability of steelmaking, steelmaking cost and the principle of smooth logistics, for steel plates that require -40℃ toughness and can be welded with ultra-large heat input, the S content needs to be controlled at ≤0.0030%.

[0039] The purpose of adding a trace amount of Nb element in steel is to effectively implement the TMCP process. When the Nb addition is less than 0.008%, the role of non-recrystallization controlled rolling cannot be effectively exerted; when the Nb addition exceeds 0.022%, it is easy to induce the formation of abnormal structures such as coarse ferrite side plate strips (FSP), Widmanstatten structure (WF), and upper bainite (Bu) under ultra-high heat input welding conditions, and promote the formation of blocky MA islands, which seriously damages the low-temperature toughness of the heat affected zone (HAZ) of ultra-high heat input welding. Therefore, the Nb content is controlled between 0.008% and 0.022%, so as to obtain the best TMCP effect without damaging the low-temperature toughness of the HAZ of ultra-high heat input welding.

[0040] The purpose of adding a trace amount of Ti to steel is to combine with N in steel to generate highly stable TiN particles, control the growth of austenite grains in the welding HAZ area, inhibit the generation of abnormal organizations such as coarse ferrite side plate strips (FSP), Widmanstatten structure (WF), upper bainite (Bu), etc., and improve the weldability of steel and the low-temperature toughness of HAZ. The Ti content added in the steel should match the N content in the steel. The matching principle is that TiN cannot precipitate in the liquid steel but must precipitate in the solid phase; when the Ti content is too little (<0.008%), the number of TiN particles formed is insufficient to inhibit the growth of austenite grains in HAZ, inhibit the generation of coarse FSP, Bu and other brittle organizations, and improve the low-temperature toughness of HAZ; when the Ti content is too much (>0.016%), large-sized TiN particles may be precipitated. Such large-sized TiN particles not only cannot inhibit the growth of austenite grains in HAZ, but become the starting point of crack initiation; therefore, the optimal control range of Ti content is 0.008% to 0.016%.

[0041] In order to ensure that during the cooling process of the welding thermal cycle, Al and N are inhibited from combining to form AlN, B and N are promoted to combine to form BN, and the low-temperature toughness of the HAZ of the steel plate in ultra-large heat input welding is improved, the lower the Als content in the steel, the better; however, as the Als content in the steel decreases to below 0.005%, the molten steel is not fully deoxidized, the inclusion content in the steel increases dramatically, and the low-temperature toughness, impact resistance and weldability of the steel plate are seriously deteriorated; however, when the Als content in the steel exceeds 0.015%, the precipitation of BN is inhibited during the ultra-large heat input welding thermal cycle, and a sufficient amount of BN cannot be formed; in summary, the appropriate Als content is controlled at 0.005% to 0.015%.

[0042] The control range of N corresponds to the control range of Ti and B. For ultra-high heat input welded steel plates, the best Ti / N is between 1.0 and 2.0. If the N content is too low, the number of TiN particles generated is small and the size is large, which cannot improve the weldability of the steel, but is harmful to the weldability; however, when the N content is too high, the free [N] in the steel increases, especially under ultra-high line energy welding conditions, the free [N] content in the heat affected zone (HAZ) increases sharply, which seriously damages the low temperature toughness of the HAZ and deteriorates the weldability of the steel. Therefore, the N content is controlled at 0.0055% to 0.0085%.

[0043] The B element is the most critical alloy element for improving the ultra-high heat input welding of steel plates. B has a dual role in the ultra-high heat input welding process. First, it can be dissolved in B and segregated on the original austenite grain boundaries in the HAZ, thereby inhibiting the nucleation of grain boundary ferrite and forming abnormal structures such as coarse ferrite side plate strips (FSP), Widmanstätten structure (WF), and upper bainite (Bu); second, the BN particles in the austenite grains are used as heterogeneous nucleation points to promote the nucleation of equiaxed ferrite in the HAZ austenite grains and refine the microstructure of the HAZ in ultra-high heat input welding. In order to achieve the above purposes, the B content in the steel is controlled at 0.0012% to 0.0030%.

[0044] Ca treatment of steel can, on the one hand, further purify the molten steel, and on the other hand, modify the sulfides in the steel to make them non-deformable, stable and fine spherical sulfides, inhibit the hot brittleness of S, improve the low-temperature toughness and Z-direction properties of the steel, and improve the anisotropy of the toughness of the steel plate; the amount of Ca added depends on the S content in the steel. If the Ca addition is too low, the treatment effect is not great; if the Ca addition is too high, the size of the Ca(O,S) formed is too large, the number of inclusions in the steel increases, the brittleness also increases, and it can become the starting point of the fracture crack, reducing the low-temperature toughness and weldability of the steel, while also reducing the purity of the steel and polluting the molten steel; therefore, the appropriate range of Ca content is 0.0010% to 0.0035%.

[0045] Solidification co-segregation index (%C) 0.5 ×[10.2(%P)+30.7(%S)+3.3(%Mn)]×[7.8(%Si)+5.6(%Als)]≤2.52.

[0046] The internal segregation of the slab has a great influence on the impact resistance of the steel plate. Since the structure of the segregated area of ​​the steel plate is different from that of the non-segregated area, the segregated area not only has abnormal brittle structure (such as high-carbon bainite, high-carbon martensite and M / A island, etc.), but also is rich in a large number of harmful inclusions (such as elongated MnS inclusions, Al 2 O 3 and SiO 2Inclusions such as C, Mn, P, S, etc., seriously deteriorate the impact resistance of the steel plate. It is extremely important to inhibit the co-segregation of C, Mn, P, S and the promotion of Si and Als to the co-segregation of C, Mn, P, S, reduce the intrinsic segregation during the solidification of molten steel, and ensure the impact resistance of the steel plate; C, Mn, P, S, Si and Als can inhibit the co-segregation of C, Mn, P, S and the promotion of Si and Als to the co-segregation of C, Mn, P, S if they meet the above formula; this is one of the key technologies of the present invention.

[0047] [1.27(%B)] / [(%N)-0.292(%Ti)-0.518(%Als)-0.147(%Nb)]≤61.5;

[0048] 1) Ensure that B and N form BN particles during heating, rolling and accelerated cooling, eliminate the existence of solid solution [B] in the steel, eliminate the phase transformation strengthening effect of B, and achieve high toughness, excellent impact resistance and toughness / plasticity matching of the steel plate;

[0049] 2) Both Ti and B are strong nitride-forming elements, but during the welding thermal cycle, TiN and BN have fundamentally different effects. TiN has high stability at relatively high temperatures (≤1250℃) and can effectively inhibit the growth of HAZ grains far from the fusion line. However, near the fusion line (FL), the peak temperature of the welding thermal cycle becomes very high (>1250℃), and the TiN particles dissolve, which cannot effectively inhibit the growth of HAZ grains. Although BN particles are not as stable as TiN particles at high temperatures and are completely dissolved in the steel during the welding thermal cycle heating and insulation process, due to the high diffusivity of B in steel, BN re-precipitates quickly during the welding thermal cycle cooling process (the diffusivity of Ti and Als is very low, and TiN and AlN cannot re-precipitate during the welding thermal cycle cooling process, that is, the precipitation kinetics of TiN and AlN are very slow). Due to the special crystal structure, BN particles have a low-energy orientation relationship with ferrite (100) BN / / (110) α BN particles can become an effective position for ferrite nucleation, promote the nucleation of equiaxed ferrite grains in austenite crystals, refine the HAZ structure, and improve the low-temperature toughness and crack resistance and crack arrest properties of the HAZ in ultra-large heat input welding heat affected zone; this is one of the key technologies of the present invention.

[0050] Ti / N≤2.0;

[0051] Ensure that [N] is excessive during the reaction of [Ti]+[N]→TiN: 1) Inhibit Oswald ripening of TiN particles and form fine dispersed TiN particles; 2) There is enough residual solid solution [N] to combine with [B], eliminate the phase transformation strengthening effect of [B] in the TMCP process, and ensure that the steel plate has excellent strength and toughness / strong plasticity matching and impact resistance; 3) Eliminate excess solid solution [N] in the welding heat affected zone, that is, the solid solution [N] is completely eaten by the nitrogen-fixing elements Ti, Als, B and Nb, eliminate the solid solution [N] in the HAZ of the steel plate welding, eliminate the huge embrittlement effect of the solid solution [N] on the ferrite structure, and improve the low-temperature toughness and crack resistance and crack arrest properties of the HAZ; this is one of the key technologies of the present invention.

[0052] Ca treatment, Ca / S ratio controlled at 1-3; 1.0×10 －3 ≤(%Ca)×(%S) 0.28 ≤3.0×10 －3 .

[0053] While ensuring sulfide spheroidization and minimizing the influence of inclusions on low-temperature toughness and weldability, a certain number of Ca(O,S) particles are evenly and finely distributed in the steel to inhibit the excessive growth of austenite grains in the heat-affected zone of high heat input welding; secondly, Ca treatment is used to form a certain number of ultra-fine Ca(O,S) inclusions to promote the precipitation and growth of BN particles in the Ca(O,S) inclusions, that is, BN particles often nucleate and grow non-uniformly on the Ca(O,S) inclusions, and then the BN particles promote the nucleation of ferrite in the austenite grains in the heat-affected zone, divide the austenite grains in the heat-affected zone, refine the grain size of the heat-affected zone, and improve the low-temperature toughness of the heat-affected zone of welding.

[0054] The steel plate structure of the present invention with extremely excellent ultra-low temperature impact toughness (-50°C) is uniform and fine equiaxed ferrite + a small amount of dispersed bainite.

[0055] The method for manufacturing the steel plate with high toughness, high ductility and ultra-large heat input having strong impact resistance according to the present invention comprises the following steps:

[0056] 1) Smelting and casting

[0057] Smelting and casting into slabs according to the above composition; wherein the casting adopts continuous casting, the superheat is controlled at 10-30°C, and the continuous casting soft pressure is controlled at 2-4%;

[0058] 2) Slab heating

[0059] The slab heating temperature is controlled between 1080 and 1150°C;

[0060] 3) Controlled rolling

[0061] The start rolling temperature of non-recrystallization controlled rolling is 810-920℃, the rolling pass reduction rate is ≥7%, the cumulative reduction rate is ≥50%, and the final rolling temperature is 800-850℃;

[0062] 4) Controlled cooling

[0063] After the controlled rolling, the steel plate is immediately subjected to the first stage of accelerated cooling, with the steel plate starting cooling temperature at 780-840°C, the steel plate stopping cooling temperature at 700-750°C, and the cooling rate at 1-4°C / s;

[0064] Then the steel plate is naturally air-cooled to 620-680°C and the second stage of accelerated cooling is carried out. The cooling temperature of the steel plate is 300-400°C and the cooling rate is ≥5°C / s. Then the steel plate is naturally air-cooled to room temperature.

[0065] In addition to meeting the above requirements, the following relationships must also be met:

[0066] [ξ×(T 开冷2 )×(DI OL )×(V c2 )] / {H×(T 开轧 )×[Ar 3 -(T 开冷2 )]×(T 停冷2 )}≥7.6×10 -3 ;in,

[0067] ξ is the cumulative reduction rate without recrystallization, unit: %;

[0068] T 开轧 is the rolling temperature before recrystallization, unit: °C;

[0069] T 开冷2 is the starting temperature of the second stage accelerated cooling, in °C;

[0070] V c2 is the accelerated cooling rate in the second stage, in °C / s;

[0071] T 停冷2 The second stage accelerated cooling stop temperature, in °C;

[0072] DI OL is the online hardenability index,

[0073] DI OL =0.51(%C) 0.5 [1+0.7(%Si)][1+3.33(%Mn)][1+0.35(%Cu)][1

[0074] +0.36(%Ni)][1+2.16(%Cr)][1+3(%Mo)][1+1.75(%V)][1+

[0075] 1.77(%Al)][1+200(%B)]×25.4, unit: mm;

[0076] H is the thickness of the finished steel plate, in mm;

[0077] Ar 3 is the starting temperature of phase transformation from austenite to ferrite under deformation, in °C;

[0078] Ar 3 =910-310(%C)-80(%Mn)-20(%Cu)-15(%Cr)-55(%Ni)

[0079] -80(%Mo)-0.35(t-8), t is the thickness of the finished steel plate, unit: mm.

[0080] In the method for manufacturing the steel plate of the present invention:

[0081] In order to ensure the excellent "three properties" (soundness, homogeneity and purity) of the slab internal quality, the continuous casting superheat is controlled between 10 and 30°C, and the continuous casting light pressure is controlled between 2% and 4%.

[0082] In order to ensure that the initial austenite grains are uniform and fine, the slab low-temperature heating process must be adopted while ensuring the complete solid solution of the micro-alloying element Nb. The slab heating temperature is controlled between 1080 and 1150°C to ensure that the austenite grain size of the original slab is uniform and relatively fine.

[0083] The starting rolling temperature of non-recrystallization controlled rolling is 810-920℃, the rolling reduction rate per pass is ≥7%, the cumulative reduction rate is ≥50%, and the final rolling temperature is 800-850℃.

[0084] After the controlled rolling is completed, the steel plate immediately undergoes the first stage of accelerated cooling. The steel plate starts cooling at 780-840°C, stops cooling at 700-750°C, and the cooling rate is 1°C / s≤≤4°C / s. The steel plate is then naturally air-cooled to 620-680°C and undergoes the second stage of accelerated cooling. The steel plate stops cooling at 300-400°C, and the cooling rate is ≥5°C / s. The steel plate is then naturally air-cooled to room temperature.

[0085] In addition to meeting the above requirements, the following relationships must also be met:

[0086] [ξ×(T 开冷2 )×(DI OL )×(V c2 )] / {H×(T 开轧 )×[Ar 3 -(T 开冷2 )]×(T 停冷2 )}≥7.6×10 -3 .

[0087] As the cumulative reduction rate ξ of non-recrystallization increases, the rolling temperature T 开轧 The flattening degree of austenite is enhanced, the density of crystal defects such as dislocation, deformation band and grain boundary step in austenite is increased, which causes the density of ferrite nucleation sites in austenite to increase, refines the TMCP steel plate structure, and improves the low-temperature toughness of the steel plate.

[0088] With the second stage accelerated cooling, the cooling temperature T 开冷2 When the temperature decreases within a certain range (620-680℃), the amount of proeutectoid ferrite increases, the carbon content of the untransformed austenite increases, and after accelerated cooling, the amount of ferrite in the microstructure of the steel plate increases and the amount of bainite decreases. Bainite is more easily embedded in the ferrite body and the hardness of bainite increases. It is easier to realize the Oroman dislocation bypass mechanism during the deformation process, and good uniform elongation and fracture elongation (i.e., impact resistance) can be obtained;

[0089] With the second stage accelerated cooling stop temperature T 停冷2 The thickness H of the finished steel plate decreases, and the strength of the steel plate increases. As the alloy content of the steel plate increases (i.e., Pcm) and the second stage accelerated cooling rate V c2 Increases, and the strength of the steel plate increases.

[0090] Comprehensive steel plate strength, toughness, impact resistance and ultra-high heat input weldability and composition, thickness, cumulative reduction rate without recrystallization, rolling temperature, second stage accelerated cooling start temperature and stop temperature and second stage accelerated cooling rate V c2 The above relationship is formed to ensure the strength, low temperature toughness, impact resistance and ultra-large heat input weldability of the steel plate. This is one of the key technologies of the present invention.

[0091] Beneficial effects of the present invention:

[0092] The present invention does not add precious alloy elements Cu and Ni, designs low-cost component combinations, and matches the corresponding TMCP process to successfully mass-produce 500MPa-grade steel plates with impact resistance and ultra-large heat input welding capabilities; the steel plates of the present invention adopt online controlled TMCP process manufacturing technology, which not only greatly reduces the manufacturing cost of the entire process and shortens the manufacturing cycle of the steel plates, but also creates huge value for the enterprise and realizes the green and environmental protection of the manufacturing process. The high performance and high added value of the steel plate are mainly reflected in its high strength, high toughness, excellent impact resistance and crack resistance and arrest properties. At the same time, the steel plate's ultra-large heat input welding performance is also excellent. It has successfully solved the key technical problems that low C content, low carbon equivalent, low cost (that is, it does not contain precious alloy elements such as Cu and Ni) and high strength, high toughness, and excellent impact resistance are conflicting and difficult to reconcile in composition design and process design, greatly improving the safety and stability of large heavy steel structures; ultra-large heat input welding saves users the cost of steel component manufacturing (ultra-large heat input high-efficiency welding), greatly shortens the time of steel component manufacturing for users, and creates huge value for users. Therefore, this type of steel plate is not only a high value-added, green and environmentally friendly product.

[0093] The manufacturing technology of the steel plate of the present invention can not only reduce the manufacturing cost of the steel plate and shorten the production cycle of the steel plate; secondly, the excellent weldability of the steel plate and the ultra-large heat input welding greatly improve the user's on-site welding efficiency, save the user's component manufacturing cost, shorten the user's component manufacturing time, and create huge value for the user; more importantly, the steel plate has excellent impact resistance, ensuring the safety and reliability of the steel structure throughout its service life cycle, so this type of steel plate is a high value-added, green and environmentally friendly product. BRIEF DESCRIPTION OF THE DRAWINGS

[0094] Figure 1 This is a microstructure photograph of the steel in Example 5 of the present invention. DETAILED DESCRIPTION

[0095] The present invention will be further described below in conjunction with the embodiments and drawings.

[0096] The composition of the steel plate of the embodiment of the present invention is shown in Table 1, and the remainder includes Fe and unavoidable impurities. Tables 2 to 4 are the manufacturing process parameters of the steel of the embodiment of the present invention, and Table 5 is the performance parameters of the steel of the embodiment of the present invention.

[0097] Depend on Figure 1 It can be seen that the microstructure of the steel plate is fine ferrite + a small amount of dispersed bainite, the average grain size is less than 20μm, and the cold deformation morphology of the proeutectoid ferrite is clearly visible.

[0098] In summary, the present invention does not add precious alloy elements Cu and Ni, combines component design at low cost, and matches the corresponding special TMCP process, fully taps the matching potential between component design and TMCP process, and obtains YP355MPa grade excellent impact resistance steel plate with excellent mechanical properties and weldability.

[0099] The steel plate manufacturing technology of the present invention not only greatly reduces the manufacturing cost of the entire process and shortens the manufacturing cycle of the steel plate, but also creates huge value for the enterprise and realizes the green and environmental protection of the manufacturing process. The high performance and high added value of the steel plate are mainly reflected in the high strength, excellent low temperature toughness and impact resistance. At the same time, the weldability of the steel plate (especially the ultra-large heat input weldability) is also excellent. It successfully solves the key technical problems that low C content, low carbon equivalent, low cost (i.e., no or trace amounts of precious alloy elements such as Cu and Ni) and high strength, impact resistance and ultra-large heat input weldability conflict with each other in composition design and process design and are difficult to reconcile, and greatly improves the safety and stability of large heavy steel structures.

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Claims

1. High toughness, high ductility and ultra-large heat input steel plate with strong impact resistance, the weight percentage of its components is: C：0.09～0.15％ Si: ≤0.20% Mn: 1.30%~1.60% P：≤0.013％ S：≤0.0030％ Nb: 0.008%~0.022% Ti: 0.008%~0.016% Als: 0.005%~0.015% N：0.0055％～0.0085％ B：0.0012％～0.0030％ Ca: 0.0010%~0.0035% The balance includes Fe and other unavoidable inclusions; and the contents of the above elements must simultaneously satisfy the following relationship: Solidification co-segregation index: (%C) 0.5 ×[10.2(%P)+30.7(%S)+3.3(%Mn)]×[7.8(%Si)+5.6(%Als)]≤2.52; [1.27(%B)] / [(%N)-0.292(%Ti)-0.518(%Als)-0.147(%Nb)]≤61.5; Ti / N≤2.0; Ca treatment, Ca / S ratio = 1.0-3.0, 1.0×10 －3 ≤(%Ca)×(%S) 0.28 ≤3.0×10 －3 .

2. The steel plate with high toughness, high ductility and ultra-large heat input having strong impact resistance as claimed in claim 1, characterized in that: The balance is Fe and other inevitable inclusions.

3. The steel plate with high toughness, high ductility and ultra-large heat input having strong impact resistance as claimed in claim 1 or 2, characterized in that: The microstructure of the steel plate is uniform and fine equiaxed ferrite + a small amount of dispersed bainite.

4. The steel plate with high toughness, high ductility and ultra-large heat input having strong impact resistance as claimed in claim 1, 2 or 3, characterized in that: The steel plate has a yield strength of ≥355MPa, a tensile strength of ≥490MPa, a uniform elongation of ≥22%, a post-fracture elongation, i.e., an impact resistance characteristic of ≥40%, a -40°C transverse impact toughness KV2 of ≥100J, and a KV2 (-40°C) of a welding HAZ of ≥70J.

5. The method for manufacturing a steel plate having high impact resistance, high toughness, high ductility and ultra-large heat input as claimed in any one of claims 1 to 4, characterized in that: The steps include: 1) Smelting and casting Smelting and casting the composition according to claim 1 or 2 into a slab; wherein the casting is continuous casting, the superheat is controlled at 10-30°C, and the continuous casting soft pressure is controlled at 2-4%; 2) Slab heating The slab heating temperature is controlled between 1080 and 1150°C; 3) Controlled rolling The start rolling temperature of non-recrystallization controlled rolling is 810-920℃, the rolling pass reduction rate is ≥7%, the cumulative reduction rate is ≥50%, and the final rolling temperature is 800-850℃; 4) Controlled cooling After the controlled rolling, the steel plate is immediately subjected to the first stage of accelerated cooling, with the steel plate starting cooling temperature at 780-840°C, the steel plate stopping cooling temperature at 700-750°C, and the cooling rate at 1-4°C / s; Then the steel plate is naturally air-cooled to 620-680°C and the second stage of accelerated cooling is carried out. The cooling temperature of the steel plate is 300-400°C and the cooling rate is ≥5°C / s. Then the steel plate is naturally air-cooled to room temperature. In addition to meeting the above requirements, the following relationships must also be met: [ξ×(T 开冷2 )×(DI OL )×(V c2 )] / {H×(T 开轧 )×[Ar3-(T 开冷2 )]×(T 停冷2 )}≥7.6×10 -3 ; among them, ξ is the cumulative reduction rate without recrystallization, unit: %; T 开轧 is the rolling temperature before recrystallization, unit: °C; T 开冷2 is the starting temperature of the second stage accelerated cooling, in °C; V c2 is the accelerated cooling rate in the second stage, in °C / s; T 停冷2 The second stage accelerated cooling stop temperature, in °C; DI OL is the online hardenability index, DI OL =0.51(%C) 0.5 [1+0.7(%Si)][1+3.33(%Mn)][1+0.35(%Cu)][1+0.36(%Ni)][1+2.16(%Cr)][1+3(%Mo)][1+1.75(%V)][1+1.77(%Al)][1+200(%B)]×25.4, unit: mm; H is the thickness of the finished steel plate, in mm; Ar3 is the starting temperature of phase transformation from austenite to ferrite under deformation, in °C; Ar3=910-310(%C)-80(%Mn)-20(%Cu)-15(%Cr)-55(%Ni)-80(%M o)-0.35(t-8), t is the thickness of the finished steel plate, unit: mm.

6. The method for manufacturing a steel plate having high toughness, high ductility and ultra-large heat input with strong impact resistance as claimed in claim 5, characterized in that: The microstructure of the steel plate is uniform and fine equiaxed ferrite + a small amount of dispersed bainite.

7. The method for manufacturing a steel plate having high toughness, high ductility and ultra-large heat input with strong impact resistance as claimed in claim 5 or 6, characterized in that: The steel plate has a yield strength of ≥355MPa, a tensile strength of ≥490MPa, a uniform elongation of ≥22%, and a post-break elongation, i.e., an anti-impact property of ≥40%. The transverse impact toughness at -40℃ is KV2≥100J, and the KV2 (-40℃) of the welding HAZ is ≥70J.

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