WO2022095761A1 - 一种超厚规格热轧h型钢及其生产方法 - Google Patents

一种超厚规格热轧h型钢及其生产方法 Download PDF

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WO2022095761A1
WO2022095761A1 PCT/CN2021/126546 CN2021126546W WO2022095761A1 WO 2022095761 A1 WO2022095761 A1 WO 2022095761A1 CN 2021126546 W CN2021126546 W CN 2021126546W WO 2022095761 A1 WO2022095761 A1 WO 2022095761A1
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flange
rolled
rolling
thickness
content
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French (fr)
Inventor
夏勐
吴保桥
吴湄庄
邢军
汪杰
陈辉
彦井成
黄琦
彭林
何军委
丁朝晖
沈千成
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Maanshan Iron and Steel Co Ltd
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Maanshan Iron and Steel Co Ltd
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Priority to EP21888458.3A priority Critical patent/EP4242338A4/en
Priority to JP2023540157A priority patent/JP7600409B2/ja
Priority to US18/258,780 priority patent/US12281369B2/en
Publication of WO2022095761A1 publication Critical patent/WO2022095761A1/zh
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    • C—CHEMISTRY; METALLURGY
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/0068—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for particular articles not mentioned below
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B1/00—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations
    • B21B1/08—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling structural sections, i.e. work of special cross-section, e.g. angle steel
    • B21B1/088—H- or I-sections
    • B—PERFORMING OPERATIONS; TRANSPORTING
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    • B—PERFORMING OPERATIONS; TRANSPORTING
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    • C21D2211/00—Microstructure comprising significant phases
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    • C21D2211/00—Microstructure comprising significant phases
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
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Definitions

  • the invention relates to the technical field of metal material production, in particular to an ultra-thick specification hot-rolled H-section steel and a production method thereof.
  • the ultra-thick hot-rolled H-beam is its core support member.
  • Patent document CN103987866B using Ni-Cu-B-V-Ti composition system special-shaped billet, through heating-rough-rolling-finishing rolling process, using cooling between finishing passes or rapid cooling after rolling to form bainite + ferrite/m
  • the room temperature structure of the intensities produces hot-rolled H-beams with a flange thickness of 100mm to 150mm, and the yield strength level is not less than 450MPa
  • patent document CN109715842A using Nb-V-Ti composition system (Cr, Mo, Ni, Cu can be added element) special-shaped billet, through the heating-rough-rolling-finishing rolling process, using cooling between finishing rolling passes or rapid cooling after rolling to form ferrite + martensite / austenite room temperature structure, producing a flange thickness of 40mm ⁇ 140mm hot-rolled H-beam, the yield strength grade is not less than 450MPa; the above two methods respectively stipulate that the bainite content at 1/4 of the flange thickness is not less than 60% and the ferrite
  • Patent documents CN105586534B, CN103938079B, CN110484822A, CN108893675A respectively adopt V, V-Ti, Ni-V-Ti, Ni-Nb-V-Mo component system special-shaped billet, through heating-blanking-rough rolling-air cooling/water cooling process, The room temperature structure of ferrite + pearlite/bainite is formed, and hot-rolled H-beam steel with a flange thickness of 45mm or less is produced. The above four methods do not control the cooling and organization of the core.
  • Patent document CN105586534B using Ni-V-Ti composition system special-shaped billet, through the process of heating-blanking-rough rolling-air cooling, the room temperature structure of ferrite + pearlite is formed, and the hot-rolled H-beam with flange thickness below 36mm is produced.
  • the yield strength level is 355MPa, and it has excellent low temperature resistance.
  • This method stipulates that the reduction rate of each pass in the rough rolling stage should not be less than 20%.
  • For products with a flange thickness of more than 90mm rolled from a profiled billet it is impossible to achieve, and no enhanced deformation penetration and cooling penetration are adopted. measures, the strength, plasticity, toughness and thickness direction properties of the product cannot be guaranteed.
  • Patent documents CN107964626B and CN 107747043B the former adopts Nb-B component system special-shaped billet, and forms a room temperature structure of tempered sorbite + ferrite + dispersed carbide through heating-rough-rolling-finishing-quenching-tempering process , the latter uses V-Ti-Ni-Mo-Cu-Cr-Al composition system special-shaped billet, through heating-rough rolling-finishing-quenching-off-line tempering process, to form a tempered martensite structure, which can produce yielding Hot-rolled H-beams with strength levels of 500MPa to 650MPa.
  • the above two methods are aimed at products with thin flange thickness, which need to achieve rapid cooling conditions at the full thickness.
  • the purpose of the present invention is to provide an ultra-thick specification hot-rolled H-section steel and a production method thereof, so as to solve the problem that the mechanical properties of the hot-rolled H-section steel with a flange thickness of 90 mm to 150 mm, especially the thickness direction properties, need to be improved.
  • the present invention adopts the following technical solutions.
  • the flange of the H-beam is 1/6 of the width and 1/4 of the thickness of the end, and the microstructure is calculated as an area percentage, including 85% to 98% of acicular ferrite, and the remaining structure is Bain Body or retained austenite, the bainite content is not more than 2%, the ferrite grain width size is not more than 40 ⁇ m, and the difference in the content of acicular ferrite in different regions along the flange thickness direction is not more than 16%.
  • the flange of the H-beam is 1/6 of the width and 1/4 of the thickness of the end, and the microstructure is calculated by area percentage, including 85% to 91% of acicular ferrite, and the remaining structure is shellfish.
  • the flange of the H-beam is 1/6 of the width and 1/4 of the thickness of the end, and the microstructure is calculated by area percentage, including 91% to 98% of acicular ferrite, and the remaining structure is shellfish.
  • the flange of the H-beam is 1/6 of the width and 1/4 of the thickness of the end, the tensile yield strength at room temperature is not less than 460 MPa, the tensile strength is not less than 540 MPa, and the elongation after fracture is not less than 24.0 %; -20 °C impact energy value is not less than 80J, and the thickness direction performance reaches Z35 level.
  • the thickness of the flange of the H-shaped steel is 90mm ⁇ 150mm.
  • a method for producing ultra-thick hot-rolled H-beam steel comprising the following steps: heating the billet, heating at a temperature of 1200°C to 1350°C, heating Time 120min ⁇ 180min; billet rolling, flange surface temperature shall not be lower than 1000°C after rolling; °C, and then enter the universal rolling mill for rolling. After the universal rolling is completed, the surface of the flange of the rolled piece is rapidly cooled to 480 °C ⁇ 530 °C at a cooling rate of 5 °C / s ⁇ 13 °C / s by water spray cooling. °C and then air-cooled.
  • the profiled billet is rolled, and the surface temperature of the flange after the rolling is completed is not lower than 1020°C.
  • This super-thick hot-rolled H-beam has a flange thickness of 90mm to 150mm, and has excellent comprehensive mechanical properties. It can meet the requirements of yield strength not less than 460MPa, tensile strength not less than 540MPa, and elongation after fracture not less than 24.0%.
  • the impact energy at -20°C is not lower than the requirement of 80J, especially the minimum performance in the thickness direction can reach the Z35 level, which can well meet the needs of heavy support structures such as high-rise buildings, large squares, and bridge structures.
  • the production method of this super-thick hot-rolled H-beam uses chemical composition control, rapid cooling before universal rolling, and segmental cooling after rolling, to form mainly acicular ferrite, and the rest are bainite or residual austenite.
  • the room temperature structure of the tensite is limited, and the content of acicular ferrite, grain size and bainite is limited, and the difference in the structure along the thickness direction of the flange is reduced.
  • the ultra-thick hot-rolled H-beam with excellent comprehensive mechanical properties has relatively low production cost and strong production achievability, and is suitable for mass production applications.
  • Fig. 1 is the typical microstructure diagram of H-beam of the present invention at room temperature
  • Figure 2 is a schematic diagram of the structure of a common H-beam, and the figure indicates the position of the flange at 1/6 of the width and 1/4 of the thickness from the end.
  • each element and the composition ratio are based on the following:
  • the lower limit is set to 0.04%; if the content exceeds 0.11%, when acicular ferrite is formed, carbides will precipitate in chains or short rods, destroying the continuity of the matrix. The plasticity, toughness and thickness direction properties are damaged, and the composition is close to the peritectic region, and cracks are easily formed at the end and inner fillet of the special-shaped billet, which adversely affects the weldability, and the upper limit is set at 0.11%.
  • Silicon (Si) A deoxidizing element in steelmaking, which increases strength and improves the fluidity of molten steel during continuous casting.
  • the lower limit is set to 0.10%; Promote the formation of martensite and austenite mixed structure, impair plasticity and toughness, and set the upper limit to 0.40%.
  • Chromium (Cr) Improve hardenability, synergize with Mn to increase the stability of supercooled austenite, promote the precipitation of acicular ferrite, improve the thickness direction properties to a certain extent, and also improve the strength.
  • the lower limit is set as 0.40%; if the content exceeds 1.00%, the enhancement of hardenability will reach saturation, which will also promote the precipitation of upper bainite, damage plasticity and toughness, and adversely affect weldability.
  • the upper limit is set to 1.00%.
  • the lower limit is set to 0.10%; if the content exceeds 0.40%, liquid segregation defects will be formed on the surface of the billet, and the upper limit is set to 0.40%.
  • Niobium (Nb) Precipitation during rolling, the pseudo-austenite grains grow, the critical temperature of the austenite unrecrystallized zone is increased, the strain accumulation is increased, and the acicular ferrite is refined. Improve the degree of work hardening of the surface and superficial areas, enhance the deformation penetration, and improve the plasticity, toughness and thickness direction properties. In order to obtain this effect, the lower limit is set to 0.020%; if the content exceeds 0.060%, the critical temperature of non-recrystallization of austenite is increased. The effect of ⁇ reaches saturation, the precipitates will aggregate and coarsen, reducing the pinning effect, and the upper limit is set to 0.060%.
  • Vanadium (V) Disperses and precipitates after rolling to increase the strength.
  • the lower limit is set to 0.040%; if the content exceeds 0.100%, the precipitates will be severely coarsened, and cracks are easily formed at the interface between the large particles and the matrix, which damages the plasticity. and toughness, and adversely affect weldability, set an upper limit of 0.100%.
  • Titanium (Ti) Precipitates during the heating stage and during rolling, and the pseudo-austenite grains grow too much. In order to obtain this effect, the lower limit is set to 0.010%; if the content exceeds 0.025%, the precipitates will aggregate and coarsen. The pinning effect is reduced, and the brittle failure base point is formed to impair the toughness, and the upper limit is set to 0.025%.
  • Aluminum (Al) A deoxidizing element in steelmaking, which precipitates during rolling, and the austenite grains are expected to grow excessively.
  • the lower limit is set to 0.010%; if the content exceeds 0.030%, brittle inclusions are easily formed, It will damage the plasticity, toughness and thickness direction properties, and it is easy to form nodules during the continuous casting process to cause steel breakout, and the upper limit is set at 0.030%.
  • N Nitrogen
  • the precipitation of Ti, V, and Nb requires the synergy of N elements, which significantly affects the precipitation quantity and distribution of Ti and V. With the increase of N content, the precipitation ratio increases greatly. In order to obtain this effect, the lower limit is set to 0.0060% ; If the content exceeds 0.0120%, the promotion of precipitation will reach saturation, and it will also promote the formation of island martensite, which will damage plasticity and toughness, and the upper limit is set at 0.0120%.
  • the lower limit of Nb+V+Ti is set to 0.090%; if the total content of the three elements is higher than 0.170%, the precipitation particles will coarsen seriously and damage toughness and plasticity.
  • the upper limit of Nb+V+Ti is set to 0.170%. Since Ti and V need to synergize with N element precipitation, if the ratio of the sum of V, Ti element content to N element content is less than 6.5, the N content required for precipitation will exceed the required amount of N, the steel gas content will increase, and the toughness will be damaged.
  • V+ The lower limit of Ti)/N is 6.5; if the ratio is higher than 10.5, the ratio of the precipitation of Ti and V elements to the total content is low, and the precipitation strengthening effect is insufficient, and the upper limit of (V+Ti)/N is set to 10.5.
  • the carbon equivalent CEV is a value calculated based on the above element content, and is also a reference index for evaluating weldability.
  • the CEV should not be lower than 0.30%, and the lower limit should be set at 0.30%; as the CEV increases, the pre-weld preparation workload and the post-weld cold cracking sensitivity increase.
  • the upper limit is set to 0.48%.
  • Phosphorus (P) impurity element, easy to solidify segregation and enrichment, damage plastic toughness, and adversely affect weldability, the upper limit is set at 0.015%.
  • S Sulfur
  • impurity element impurity element, rolling to form elongated inclusions, the atomic arrangement of the contact surface is disordered, the energy is high, cracks are easily generated, and the toughness and thickness direction properties are damaged.
  • the upper limit is set at 0.005%.
  • Oxygen (O) an impurity element, which forms oxide inclusions with various elements, forms a brittle failure base point, and impairs plasticity, toughness and thickness direction properties, and the upper limit is set to 0.0060%.
  • the ultra-thick hot-rolled H-beam of the present invention has its mechanical properties sampled at 1/6 of the width and 1/4 of the thickness from the flange to the end, so the microstructure is also taken as characterization.
  • the microstructure is calculated by area percentage, including 85% to 98% of acicular ferrite, the remaining structure is bainite or retained austenite, and the bainite content is not higher than 2%, of which the ferrite grain width
  • the size is not more than 40 ⁇ m, but the difference in the content of acicular ferrite in different regions along the flange thickness direction is not more than 16%.
  • the microstructure is calculated by area percentage, at 1/6 of the width and 1/4 of the thickness of the flange from the end, including 85% to 91% of acicular ferrite, and the remaining structure is Bain Body and retained austenite, the bainite content is not more than 2%, the ferrite grain width is not more than 20 ⁇ m, and the difference in the content of acicular ferrite in different regions along the flange thickness direction should not be more than 9%.
  • the microstructure is calculated by area percentage, and at the position of 1/6 of the width and 1/4 of the thickness of the flange from the end, the surface of the flange contains 91% to 98% of acicular ferrite, and the remaining structure It is bainite and retained austenite, and the bainite content is not more than 1%, of which the ferrite grain width is 20 ⁇ m to 40 ⁇ m, and the content of acicular ferrite in different regions along the flange thickness varies by 9% to 9%. 16%.
  • the ferrite content is less than 85% at 1/6 of the width and 1/4 of the thickness from the flange to the end, the total content of acicular ferrite in the entire thickness range is insufficient, and pearlescent may even appear.
  • the beneficial effect of acicular ferrite cannot be fully exerted, and the plasticity, toughness and thickness direction properties are damaged, and the lower limit of its content is set to 85%; due to the precipitation of acicular ferrite, C element must be enriched in other areas. , bainite or retained austenite is formed in the post-rolling cooling process, which cannot be completely transformed into acicular ferrite, and the upper limit of its content is set to 98%.
  • the acicular ferrite grains are in the shape of short rods, and the aspect ratio is usually 2:1 to 5:1. It is a feasible method to limit the width size. Reducing the grain size can improve the strength, plasticity and toughness, and improve the thickness direction. The uniformity of grain size is conducive to improving the performance in the thickness direction. If the width is greater than 40 ⁇ m, the comprehensive mechanical properties will be reduced, and the upper limit of the width size is set to 40 ⁇ m.
  • the content of acicular ferrite in the thickness direction of the flange differs by more than 16%, the plasticity gap between different regions will increase.
  • the crack nucleation work is easy to become the source of the crack and damage the thickness directionality.
  • the production method of the super-thick specification hot-rolled H-beam of the present invention is: molten iron pretreatment ⁇ converter smelting ⁇ argon blowing, LF furnace refining ⁇ billet heating ⁇ billet rolling ⁇ universal rolling (rapid cooling before rolling) ⁇ Segment cooling after rolling (rapid cooling + air cooling).
  • the production method of the super-thick specification hot-rolled H-section steel of the present invention adopts special-shaped billet rolling, the heating temperature of the billet is controlled at 1200°C to 1350°C, and the heating time is not less than 120 minutes.
  • the heating temperature is 1200°C ⁇ 1260°C
  • the heating time is 122min ⁇ 144min.
  • the heating temperature is 1260°C-1350°C, and the heating time is 168min-173min.
  • the purpose of heating the billet is to make the alloy composition solid solution, homogenize the structure, and reduce the rolling deformation resistance.
  • the alloy elements will not have sufficient time for solid solution, and the precipitates containing Ti and Nb will form particles with uneven size during the precipitation process, which cannot be dispersed and distributed, and cannot play the role of pinning and strengthening.
  • the lower limit is 1200 °C; if the temperature exceeds 1350 °C, the original grain size will increase, which is not conducive to the dispersion and distribution of precipitates, and it is easy to overfire to form surface and shallow surface cracks, and the upper limit is set at 1350 °C.
  • the heating time is less than 120min, the core of the billet cannot be burned through, and the solid solution and homogenization of alloy elements are not sufficient. more than 180min.
  • the surface temperature of the flange is not lower than 1000°C.
  • the surface temperature of the flange is not lower than 1020°C.
  • the surface temperature of the flange is not lower than 1050°C.
  • the purpose of billet rolling is to shape the billet and provide a suitable billet shape for universal rolling.
  • the purpose of controlling the surface temperature of the flange after the billet rolling is completed is to form a certain temperature gradient from the surface and the core through rapid cooling to enhance the deformation penetration in the universal rolling stage.
  • the flange core temperature is higher than the surface temperature. If the surface temperature is lower than 1000 °C, the core temperature is lower than 1100 °C.
  • the overall thermal If the capacity is small, the temperature of the core drops quickly, and an effective temperature gradient cannot be formed from the surface to the core, which affects the effect of deformation and penetration.
  • the lower limit is set at 1100 °C.
  • the surface of the flange is cooled to 700°C to 800°C at a cooling rate of not less than 20°C/s by means of water spray cooling, Then enter the universal rolling mill for rolling.
  • the flange surface is cooled to 740°C-800°C by water spray cooling at a cooling rate of not less than 22°C/s, and then enters the universal mill for rolling.
  • the flange surface is cooled to 700°C to 740°C by means of water spray cooling at a cooling rate of not less than 32°C/s, and then enters the universal mill for rolling.
  • the purpose of universal rolling is to compress the flange and web in the thickness direction to obtain the shape and size of the finished product.
  • the purpose of rapid cooling before universal rolling is to form a certain temperature gradient from the surface to the core in the thickness direction of the flange.
  • the surface temperature is low and the deformation resistance is large.
  • the deformation will gradually penetrate into the core with higher temperature and less deformation resistance.
  • Practice shows that with the increase of temperature gradient, the deformation penetration effect is enhanced, the strain accumulation of the core increases, and the difference of strain accumulation from the surface to the core decreases accordingly.
  • the easily measurable surface temperature is used as the process parameter.
  • the cooling rate is lower than 20°C/s, the cooling rate of the surface is too slow, and the heat in the core has sufficient time to conduct to the surface layer, and an effective temperature gradient cannot be obtained.
  • the cooling temperature of the flange surface is cooled to below 700°C, the temperature of the core will be low, the deformation resistance of this area will increase, which will affect the deformation penetration effect, and the energy consumption will be too large.
  • the lower limit is set to 700°C; if the temperature is higher than 800 °C, the degree of surface work hardening is insufficient, and the deformation is still concentrated on the surface, which affects the deformation penetration effect, and the upper limit is set at 800 °C.
  • the surface of the flange of the rolled piece is cooled to 480° C. at a cooling rate of 5° C./s to 13° C./s by means of water spray cooling. ⁇ 530°C, and then air-cooled, generally a cooling bed can be used.
  • the surface of the flange is cooled to 505°C to 530°C by water spray cooling at a cooling rate of 5°C/s to 9°C/s, and then air-cooled.
  • the surface of the flange is cooled to 480°C to 505°C by means of water spray cooling at a cooling rate of 9°C/s to 13°C/s, and then air-cooled.
  • the purpose of rapid cooling after universal rolling is to simulate the precipitation of massive ferrite and pearlite, while avoiding the precipitation of bainite and promoting the formation of fine acicular ferrite as much as possible.
  • the cooling rate is higher than the critical cooling rate of the two, and the upper limit is set, and the surface of the flange is cooled by water spray, giving the core sufficient cooling time through heat conduction.
  • the full thickness of the flange is controlled within the temperature range of 480°C to 580°C, and fine acicular ferrite is fully formed in the air cooling stage.
  • the cooling rate of the flange surface is lower than 5°C/s, the cooling rate of the core will decrease, and block ferrite or pearlite will be precipitated in a strip-like distribution, which will damage the toughness and plasticity.
  • the lower limit is set to 5°C/s ; If the cooling rate is higher than 13°C/s, the total cooling time will be insufficient, the heat conduction of the core will be insufficient, and the initial temperature of air cooling will be high, which will form a large amount of pearlite, which will damage the toughness and plasticity.
  • the upper limit is set to 13°C/s.
  • the cooling temperature of the flange surface is lower than 480°C, the surface and nearby areas will fall into the upper bainite precipitation range, and more than 3% bainite will be formed, while the core area is shaped like acicular ferrite, and the microstructure difference If the cooling temperature is higher than 580 °C, the initial temperature of air cooling in the core will increase, and a large amount of massive ferrite will be precipitated, which will form on the surface and nearby areas. Large acicular ferrite impairs toughness and thickness direction properties, and the upper limit is set to 580°C.
  • the ultra-thick specification hot-rolled H-section steel of the present invention has a flange thickness range of 90 mm to 150 mm, and a web thickness range of 50 mm to 120 mm at this time.
  • the thickness of the flange ranges from 90 mm to 115 mm.
  • the thickness of the flange ranges from 115 mm to 150 mm.
  • the thickness of the hot-rolled H-beam is required to be not less than 90mm, and the lower limit is set to 90mm;
  • the equipment investment is large, the production difficulty is high, and the flange is too thick, the rolling deformation penetration and the controlled cooling penetration are limited, and the upper limit is set at 150mm.
  • the flange thickness is in the range of 90mm to 150mm, considering the structural stability and production feasibility, the web thickness is 50mm to 120mm.
  • Tables 1-4 are respectively the chemical components, production process parameters, microstructure and mechanical properties of Examples 1-10 provided by the present invention:
  • the ultra-thick specification hot-rolled H-beam of the present invention is sampled on the flange at 1/6 of the width and 1/4 of the thickness from the end; according to the standard GB/T 228.1,
  • the measured tensile yield strength at room temperature should not be less than 460MPa, the tensile strength should not be less than 540MPa, and the elongation after fracture should not be less than 24.0%; according to the standard GB/T 229, the measured impact energy at -20°C
  • the value should not be less than 80J; according to the standard GB/T 5313, the measured thickness direction performance should reach the Z35 level.
  • the method of the present invention is used to produce ultra-thick hot-rolled H-beams with a flange thickness of 90 mm to 150 mm, the yield strength at room temperature reaches 464 MPa to 522 MPa, and the tensile strength is 597 MPa to 649 MPa.
  • the elongation is 24.0% ⁇ 32.0%, the impact energy at -20°C is 84J ⁇ 126J, and the performance in the thickness direction exceeds the requirements of Z35 grade.

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Abstract

本发明公开了一种超厚规格热轧H型钢及其生产方法，其化学成分按质量百分比计包括C：0.04～0.11，Si：0.10～0.40，Mn：0.40～1.00，Cr：0.40～1.00，Cu：0.10～0.40，Nb：0.020～0.060，V：0.040～0.100，Ti：0.010～0.025，Al：0.010～0.030，N：0.0060～0.0120，P：≤0.015，S：≤0.005，O：≤0.0060，0.090％≤Nb+V+Ti≤0.170％，6.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，0.30％≤CEV≤0.48％；本发明翼缘厚度90mm～150mm，综合力学性能优良，可以很好地满足高层建筑、大型广场、桥梁结构等的重型支撑结构件需求。

Description

一种超厚规格热轧H型钢及其生产方法 技术领域
本发明涉及金属材料生产技术领域，具体为一种超厚规格热轧H型钢及其生产方法。
背景技术
伴随着经济社会高速发展，高层建筑、大型场馆、桥梁主塔等建造，对安全性、舒适性和美观性方面的需求越来越重视，设计采用重型钢结构，而翼缘厚度90mm～150mm的超厚规格热轧H型钢是其核心支撑构件。
长期实践表明，热轧H型钢翼缘的综合力学性能弱于腹板，而标准GB/T2975也规定力学性能评价的在翼缘取样，业内通常以热轧H型钢翼缘作为描述对象。采用大厚度坯料轧制薄规格，通过大压下可以获得优良的力学性能，是行业所认同的。但是，采用该方法轧制翼缘厚度90mm～150mm的超厚规格热轧H型钢，则需要远超现有尺寸的超厚、超大坯料，必须投资建设新的连铸设备和轧钢设备，成本巨大、而且坯料的内部和表面质量控制难度极高，很难实现。从可实现性和经济性方面考虑，采用异型坯轧制，从坯料到成品在翼缘厚度方向的压下率为13％～30％，是可以接受的。
专利文件CN103987866B，采用Ni-Cu-B-V-Ti成分体系异型坯，经加热-粗轧-精轧工序，利用精轧道次间冷却或轧后快速冷却，形成贝氏体+铁素体/马氏体的室温组织，生产出翼缘厚度100mm～150mm的热轧H型钢，屈服强度级别不低于450MPa；专利文件CN109715842A，采用Nb-V-Ti成分体系(可以添加Cr、Mo、Ni、Cu元素)异型坯，经加热-粗轧-精轧工序，利用精轧道次间冷却或轧后快速冷却，形成铁素体+马氏体/奥氏体室温组织，生产出翼缘厚度40mm～140mm的热轧H型钢，屈服强度级别不低于450MPa；上述两种方法，分别规定在翼缘厚度1/4处的贝氏体含量不低于60％和铁素体含量不低于60％(晶粒尺寸不大于35μm)，但没对翼缘全厚度方向的组织形貌和含量进行调控。如该方法应用于翼缘厚度90mm以上规格产品生产时，厚度方向性能无法保证。
专利文件CN105586534B、CN103938079B、CN110484822A、CN108893675A，分别采用V、V-Ti、Ni-V-Ti、Ni-Nb-V-Mo成分体系异型坯，经加热-开坯-粗轧-空冷/水冷工序，形成铁素体+珠光体/贝氏体的室温组织，生产出翼缘厚度45mm以下的热轧H型钢，屈服强度级别355MPa～500MPa，满足-20℃～-40℃低温韧性要求。上述四种方法，没有对芯部的冷却和组织进行调控。在应用于翼缘厚度90mm以上规格产品生产，由于芯部起始冷却温度高、空冷全截面冷却速度低，易析出大尺寸块状铁素体，并呈链状或网状分布，析出相也易聚集粗化，产品的强度、塑性、韧性和厚度方向性能无法保证，且添加Ni、Mo等元素将增加合金成本。
专利文件CN105586534B，采用Ni-V-Ti成分体系异型坯，经加热-开坯-粗轧-空冷工序，形成铁素体+珠光体的室温组织，生产出翼缘厚度36mm以下的热轧H型钢，屈服强度级别355MPa，具有优良耐低温性能。该方法规定，在粗轧阶段的每道次压下率应不低于20％，对于由异型坯轧制翼缘厚度90mm以上规格产品，是无法达到的，也没有采取增强变形渗透和冷却渗透的措施，产品的强度、塑性、韧性和厚度方向性能无法保证。
专利文件CN107964626B和CN 107747043B，前者采用Nb-B成分体系异型坯，经加热-粗轧-精轧-淬火-回火工序，形成回火索氏体+铁素体+弥散分布碳化物的室温组织，后者采用V-Ti-Ni-Mo-Cu-Cr-Al成分体系异型坯，经加热-粗轧-精轧-淬火-离线回火工序，形成回火马氏体组织，可以生产出屈服强度级别500MPa～650MPa的热轧H型钢。上述两种方法针对的是翼缘厚度较薄的产品，需要在全厚度达到快速冷却条件，如该方法应用于翼缘厚度90mm以上规格产品生产，则全厚度无法全部达到淬火临界冷速，不能获得在线或离线热处理所需的原始组织，芯部冷却起始温度高而冷却速度低，不能形成索氏体/马氏体+弥散分布碳化物，产品的强度、塑性、韧性和厚度方向性能，且添加Ni、Mo元素将增加合金成本。
综上所述，急需一种超厚规格热轧H型钢及其生产方法来解决这些问题。
发明内容
1、要解决的问题
本发明的目的在于提供一种超厚规格热轧H型钢及其生产方法，以解决翼缘厚度90mm～150mm的热轧H型钢的力学性能，尤其是厚度方向性能有待提高的问题。
2、技术方案
为解决上述问题，本发明采用如下的技术方案。
一种超厚规格热轧H型钢，其化学成分按质量百分比计，包括C：0.04～0.11，Si：0.10～0.40，Mn：0.40～1.00，Cr：0.40～1.00，Cu：0.10～0.40，Nb：0.020～0.060，V：0.040～0.100，Ti：0.010～0.025，Al：0.010～0.030，N：0.0060～0.0120，P：≤0.015，S：≤0.005，O：≤0.0060，满足0.090％≤Nb+V+Ti≤0.170％和6.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.30％≤CEV≤0.48％。
可选的，H型钢的化学成分按质量百分比计，包括C：0.04～0.07，Si：0.10～0.30，Mn：0.80～1.00，Cr：0.40～0.90，Cu：0.10～0.25，Nb：0.040～0.060，V：0.040～0.080，Ti：0.010～0.015，Al：0.010～0.020，N：0.0060～0.0100，P：≤0.015，S：≤0.005，O：≤0.0060，满足0.090％≤Nb+V+Ti≤0.130％，6.5≤(V+Ti)/N≤8.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.30％≤CEV≤ 0.43％。
可选的，H型钢的化学成分按质量百分比计，包括C：0.07～0.11，Si：0.30～0.40，Mn：0.40～0.80，Cr：0.90～1.00，Cu：0.25～0.40，Nb：0.020～0.040，V：0.080～0.100，Ti：0.015～0.025，Al：0.020～0.030，N：0.0100～0.0120，P：≤0.015，S：≤0.005，O：≤0.0040，满足0.130％＜Nb+V+Ti≤0.170％，8.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.40％≤CEV≤0.48％。
优选的，H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，显微组织以面积百分比计，包含85％～98％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于2％，其中铁素体晶粒宽度尺寸不大于40μm，沿翼缘厚度方向不同区域的针状铁素体含量差距不大于16％。
可选的，H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，显微组织以面积百分比计，包含85％～91％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于2％，其中铁素体晶粒宽度尺寸不大于20μm，沿翼缘厚度方向不同区域的针状铁素体含量差距不大于9％。
可选的，H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，显微组织以面积百分比计，包含91％～98％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于1％，其中铁素体晶粒宽度尺寸20μm～40μm，沿翼缘厚度方向不同区域的针状铁素体含量差距9％～16％。
优选的，H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，室温拉伸屈服强度不低于460MPa，抗拉强度不低于540MPa，断后伸长率不低于24.0％；-20℃冲击功值不低于80J，厚度方向性能达到Z35级别。
优选的，H型钢的翼缘厚度为90mm～150mm。
本发明提供的另一技术方案：一种如权利要求1-8任意一项所述的超厚规格热轧H型钢的生产方法，包括以下步骤：坯料加热，加热温度1200℃～1350℃，加热时间120min～180min；开坯轧制，轧制完成后翼缘表面温度不低于1000℃；以不低于20℃/s的冷却速度喷水冷却，将翼缘表面快速冷却至700℃～800℃，然后进入万能轧机轧制，万能轧制完成后，先采用喷水冷却的方式，以5℃/s～13℃/s的冷却速度，将轧件翼缘表面快速冷却至480℃～530℃，然后空冷。
优选的，异型坯开坯轧制，轧制完成后翼缘表面温度不低于1020℃。
3、有益效果
相比于现有技术，本发明的有益效果为：
1、该超厚规格热轧H型钢，翼缘厚度90mm～150mm，综合力学性能优良，可以满足屈服强度不低于460MPa、抗拉强度不低于540MPa、断后伸长率不低于24.0％、-20℃冲击功不低于80J的要求，尤其是厚度方向性能最低可以达到Z35级别，可以很好地满足高层建筑、大型广场、桥梁结构等的重型支撑结构件需求。
2、该超厚规格热轧H型钢的生产方法，在万能轧制前快速冷却，可以在翼缘厚度方向从表面至芯部形成一定的温度梯度，在轧制过程中，表面温度低，变形抗力大，随着轧制压缩变形不断进行，变形将逐步渗透至温度更高、变形抗力小的芯部，随着温度梯度增大，变形渗透效果增强，芯部的应变积累增大，从表面到芯部的应变积累差距则随之减小，通过增大应变积累，增加形核位置、提升驱动力，促进针状铁素体析出并细化；降低翼缘厚度方向的应变积累差，有利于减少厚度方向不同区域组织含量的差距，提升组织均匀性。
3、该超厚规格热轧H型钢的生产方法，利用化学成分控制、万能轧制前快速冷却、轧后分段冷却，形成以针状铁素体为主，其余为贝氏体或残余奥氏体的室温组织，并限制针状铁素体含量、晶粒尺寸及贝氏体含量，减小沿翼缘厚度方向组织差异，利用组织、析出、固溶和细晶强化的复合作用，获得综合力学性能优良的超厚规格热轧H型钢，生产成本相对较低，生产可实现性强，适合大量生产应用。
附图说明
图1为本发明的H型钢在室温下的典型微观结构图；
图2为常见H型钢的结构示意图，图中标出翼缘距离端部的宽度1/6处、厚度1/4处的位置。
具体实施方式
本发明的超厚规格热轧H型钢，其化学成分按质量百分比计，C：0.04～0.11，Si：0.10～0.40，Mn：0.40～1.00，Cr：0.40～1.00，Cu：0.10～0.40，Nb：0.020～0.060，V：0.040～0.100，Ti：0.010～0.025，Al：0.010～0.030，N：0.0060～0.0120，P：≤0.015，S：≤0.005，O：≤0.0060，满足0.090％≤Nb+V+Ti≤0.170％和6.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.30％≤CEV≤0.48％。
进一步可选的，化学成分按质量百分比计，C：0.04～0.07，Si：0.10～0.30，Mn：0.80～1.00，Cr：0.40～0.90，Cu：0.10～0.25，Nb：0.040～0.060，V：0.040～0.080，Ti：0.010～0.015，Al：0.010～0.020，N：0.0060～0.0100，P：≤0.015，S：≤0.005，O：≤0.0060，满足0.090％≤Nb+V+Ti≤0.130％，6.5≤(V+Ti)/N≤8.5，其余为Fe及微量残余元素，根据 CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.30％≤CEV≤0.43％。
进一步可选的，化学成分按质量百分比计，C：0.07～0.11，Si：0.30～0.40，Mn：0.40～0.80，Cr：0.90～1.00，Cu：0.25～0.40，Nb：0.020～0.040，V：0.080～0.100，Ti：0.015～0.025，Al：0.020～0.030，N：0.0100～0.0120，P：≤0.015，S：≤0.005，O：≤0.0040，满足0.130％＜Nb+V+Ti≤0.170％，8.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.40％≤CEV≤0.48％。
具体来说，各元素的作用及成分配比(按质量百分比计)依据如下：
碳(C)：提升强度，为获得该效果，设定下限为0.04％；如含量超过0.11％，在形成针状铁素体时，碳化物将呈链条或短棒状析出，破坏基体连续性，损害塑性、韧性和厚度方向性能，且成分接近包晶区，异型坯端部和内圆角易形成裂纹，并对可焊性产生不利影响，设定上限为0.11％。
硅(Si)：炼钢脱氧元素，提升强度，改善钢液在连铸时的流动性，为获得该效果，设定下限为0.10％；如含量超过0.40％，提升强度作用达到饱和，也将促进马氏体和奥氏体混合组织形成，损害塑性和韧性，设定上限为0.40％。
锰(Mn)：提高淬透性，与Cr元素协同作用增大过冷奥氏体稳定性，促进针状铁素体析出，在一定程度上改善厚度方向性能，也能提升强度，为获得该效果，设定下限为0.40％；如含量超过1.00％，易造成宏观成分偏析，珠光体或贝氏体或残余奥氏体将呈条带状分布，破坏基体连续性，损害厚度方向性能，设定上限为1.00％。
铬(Cr)：提高淬透性，与Mn元素协同作用增大过冷奥氏体稳定性，促进针状铁素体析出，在一定程度上改善厚度方向性能也能提升强度，设定下限为0.40％；如含量超过1.00％，对淬透性的提升作用达到饱和，也将促进上贝氏体析出，损害塑性和韧性，并对可焊性产生不利影响，设定上限为1.00％。
铜(Cu)：提升强度，为获得该效果，设定下限为0.10％；如含量超过0.40％，坯料表面将形成液析缺陷，设定上限为0.40％。
铌(Nb)：在轧制期间析出，拟制奥氏体晶粒长大，提升奥氏体未再结晶区临界温度，增大应变积累，有助于细化针状铁素体，也能提升表面及浅表区加工硬化程度，增强变形渗透，改善塑性、韧性和厚度方向性能，为了获得该效果，设定下限为0.020％；如含量超过0.060％，提升奥氏体未再结晶临界温度的作用达到饱和，析出物将聚集粗化，降低钉扎作用，设定上限为0.060％。
钒(V)：在轧后弥散析出，提升强度，为获得该效果，设定下限为0.040％；如含量超过0.100％，析出物粗化严重，大颗粒与基体交界处易萌生裂纹，损害塑性和韧性，并对可焊性 产生不利影响，设定上限为0.100％。
钛(Ti)：在加热阶段和轧制期间析出，拟制奥氏体晶粒过分长大，为获得该效果，设定下限为0.010％；如含量超过0.025％，析出物将聚集粗化，降低钉扎作用，并形成脆性破坏基点进而损害韧性，设定上限为0.025％。
铝(Al)：炼钢脱氧元素，在轧制期间析出，拟制奥氏体晶粒过分长大，为获得该效果，设定下限为0.010％；如含量超过0.030％，易形成脆性夹杂，损害塑性、韧性和厚度方向性能，连铸过程中也易结瘤造成漏钢，设定上限为0.030％。
氮(N)：Ti、V、Nb析出需要N元素协同，其显著影响Ti和V的析出数量和分布，随着N含量增加，析出比例大幅增加，为获得该效果，设定下限为0.0060％；如含量超过0.0120％，对析出的促进作用达到饱和，也会促进岛状马氏体形成，损害塑性和韧性，设定上限为0.0120％。
为了充分发挥Nb、V、Ti元素析出的有益作用，设定Nb+V+Ti下限为0.090％；如三种元素的总含量高于0.170％，则析出颗粒粗化严重，损害韧性和塑性，设定Nb+V+Ti上限为0.170％。由于Ti和V需要协同N元素析出，如V、Ti元素含量总和与N元素含量的比例低于6.5，则超过析出所需的N含量，钢种气体含量增加，损害韧性，设定(V+Ti)/N下限为6.5；如比例高于10.5，则Ti、V元素的析出量占其总含量的比例较低，析出强化作用不足，设定(V+Ti)/N上限为10.5。
根据标准GB/T 1591规定，碳当量CEV是基于上述元素含量计算的数值，也是评价可焊性的参考指标。为了有效发挥各化学成分的作用，CEV不低于0.30％，设定下限为0.30％；由于随着CEV增大，在焊接使用时，焊前准备工作量、焊后冷裂敏感性均增大，为了便于产品后续焊接使用，设定上限为0.48％。
磷(P)：杂质元素，易于凝固偏析和富集，损害塑韧性，对可焊性有不利影响，设定上限为0.015％。
硫(S)：杂质元素，轧制形成长条状夹杂，接触面原子排列紊乱、能量较高，易产生裂纹，损害韧性和厚度方向性能，设定上限为0.005％。
氧(O)：杂质元素，与多种元素形成氧化物夹杂，形成脆性破坏基点，损害塑性、韧性和厚度方向性能，设定上限为0.0060％。
本发明的超厚规格热轧H型钢，根据标准GB/T 2975规定，力学性能在翼缘距离端部的宽度1/6处、厚度1/4处取样，所以显微组织也以该处作为表征。显微组织以面积百分比计，包含85％～98％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于2％，其中铁素体晶粒宽度尺寸不大于40μm，但沿翼缘厚度方向不同区域的针状铁素体含量差距 不大于16％。
进一步可选的，显微组织以面积百分比计，在翼缘距离端部的宽度1/6处、厚度1/4处，包含85％～91％的针状铁素体，剩余组织为贝氏体和残余奥氏体，贝氏体含量不高于2％，其中铁素体晶粒宽度尺寸不大于20μm，沿翼缘厚度方向不同区域的针状铁素体含量差距应不大于9％。
进一步可选的，显微组织以面积百分比计，在翼缘距离端部的宽度1/6处、厚度1/4处，翼缘表面包含91％～98％的针状铁素体，剩余组织为贝氏体和残余奥氏体，贝氏体含量不高于1％，其中铁素体晶粒宽度尺寸20μm～40μm，沿翼缘厚度方向不同区域的针状铁素体含量差距9％～16％。
由于针状铁素体的长轴方向不固定，晶界形成互锁结构，在翼缘厚度方向的分布也是同样的情况，其在提升塑性、韧性和厚度方向性能方面明显优于贝氏体和珠光体，提升强度的作用也高于珠光体，是提升产品综合力学性能的关键组织。
如在翼缘距离端部的宽度1/6处、厚度1/4处，状铁素体含量低于85％，则在全厚度范围内针状铁素体总含量不足，甚至可能会出现珠光体，针状铁素体的有益的作用不能充分发挥，损害塑性、韧性和厚度方向性能，设定其含量下限为85％；由于针状铁素体析出时，C元素必然在其他区域富集，在轧后冷却过程中形成贝氏体或残余奥氏体，不能完全转变为针状铁素体，设定其含量上限为98％。
由于析出的贝氏体分布相对集中，破坏了基体连续性，如其含量超过2％，则损害塑性和韧性，设定其含量上限为2％。
针状铁素体晶粒成短棒状，长径比通常为2:1～5:1，限定其宽度尺寸是可行的方法，减小晶粒尺寸，能够提升强度、塑性和韧性，改善厚度方向晶粒尺寸均匀性，有利于提升厚度方向性能，如宽度大于40μm，则综合力学性能降低，设定其宽度尺寸上限为40μm。
如翼缘厚度方向的针状铁素体含量相差超过16％，则不同区域的塑性差距增大，在受到沿厚度方向的拉伸作用时，不同区域间无法协调变形，界面两侧应力差超过裂纹形核功，易成为裂纹源头，损害厚度方向性，设定沿翼缘厚度方不同区域的针状铁素体含量差距不大于16％。
本发明的超厚规格热轧H型钢生产方法，主要生产工序为：铁水预处理→转炉冶炼→吹氩，LF炉精炼→坯料加热→开坯轧制→万能轧制(开轧前快速冷却)→轧后分段冷却(快速冷却+空冷)。
本发明的超厚规格热轧H型钢生产方法，采用异型坯轧制，坯料的加热温度控制在1200℃～1350℃，加热时间不低于120min。
进一步可选的，加热温度1200℃～1260℃，加热时间122min～144min。
进一步可选的，加热温度1260℃～1350℃，加热时间168min～173min。
坯料加热的目的，是让合金成分固溶、组织均匀化、降低轧制变形抗力。
如温度低于1200℃，则合金元素无充足时间固溶，含Ti、Nb元素的析出物在析出过程中形成尺寸不均匀的颗粒，也不能弥散分布，无法发挥钉扎和强化作用，设定下限为1200℃；如温度超过1350℃，则原始晶粒尺寸增大，不利于析出物弥散分布，而且易过烧形成表面和浅表层裂纹，设定上限为1350℃。
如加热时间低于120min，则坯料芯部无法烧透，合金元素固溶和均匀化不充分，设定下限为120min；从减少氧化烧损、降低加热能耗等组产经济性方面考虑，不宜超过180min。
本发明的超厚规格热轧H型钢生产方法，完成开坯轧制后，翼缘表面温度不低于1000℃。
进一步可选的，翼缘表面温度不低于1020℃。
进一步可选的，翼缘表面温度不低于1050℃。
开坯轧制的目的，是对坯料进行整形，为万能轧制提供合适的坯形。完成开坯轧制后控制翼缘表面温度目的，是为了通过快速冷却，从表面和芯部形成一定的温度梯度，增强万能轧制阶段的变形渗透。完成开坯轧制后，翼缘芯部温度高于表面温度，如表面温度低于1000℃，则芯部温度低于1100℃，在万能轧制阶段对翼缘表面进行快速冷却时，总体热容小，芯部温降快，从表面到芯部的无法形成有效的温度梯度，影响变形渗透效果，设定下限为1100℃。
本发明的超厚规格热轧H型钢生产方法，在万能轧制前，采用喷水冷却的方式，以不低于20℃/s的冷却速度，将翼缘表面冷却至700℃～800℃，然后进入万能轧机轧制。
进一步可选的，在万能轧制前，采用喷水冷却的方式，以不低于22℃/s的冷却速度，将翼缘表面冷却至740℃～800℃，然后进入万能轧机轧制。
进一步可选的，在万能轧制前，采用喷水冷却的方式，以不低于32℃/s的冷却速度，将翼缘表面冷却至700℃～740℃，然后进入万能轧机轧制。
万能轧制的目的，是对翼缘和腹板厚度方向进行压缩变形，获得成品形状和尺寸。
在万能轧制前快速冷却的目的，是在翼缘厚度方向从表面至芯部形成一定的温度梯度，在轧制过程中，表面温度低，变形抗力大，随着轧制压缩变形不断进行，变形将逐步渗透至温度更高、变形抗力小的芯部。实践表明，随着温度梯度增大，变形渗透效果增强，芯部的应变积累增大，从表面到芯部的应变积累差距则随之减小。通过增大应变积累，增加形核位置、提升驱动力，促进针状铁素体析出并细化；降低翼缘厚度方向的应变积累差，有利于减少厚度方向不同区域组织含量的差距，提升组织均匀性。
由于生产过程中不便于快速测定翼缘芯部的温度，以易于测得的表面温度作为工艺参数。
如冷却速度低于20℃/s，则表面冷速过慢，芯部热量有充足的时间传导至表层，无法获得有效的温度梯度。
如翼缘表层冷却温度冷却至700℃以下，则芯部的温度较低，该区域变形抗力增大，影响变形渗透效果，而且能源消耗过大，设定下限为700℃；如温度高于800℃，则表面加工硬化程度不足，变形依然集中在表面，影响变形渗透效果，设定上限为800℃。
本发明的超厚规格热轧H型钢生产方法，完成万能轧制后，采用喷水冷却的方式，以5℃/s～13℃/s的冷却速度，将轧件翼缘表面冷却至480℃～530℃，然后空冷，一般可以采用冷床。
进一步可选的，完成万能轧制后，采用喷水冷却的方式，以5℃/s～9℃/s的冷却速度，将翼缘表面冷却至505℃～530℃，然后空冷。
进一步可选的，完成万能轧制后，采用喷水冷却的方式，以9℃/s～13℃/s的冷却速度，将翼缘表面冷却至480℃～505℃，然后空冷。
完成万能轧制后快速冷却的目的，是拟制块状铁素体和珠光体析出，同时避免贝氏体析出，尽可能多得促进细小的针状铁素体形成。快速通过先共析铁素体和珠光体析出温度区间，冷却速度高于两者的临界冷速，并设定上限，对翼缘表面进行喷水冷却，给予芯部通过热传导充分的冷却时间。终冷时，翼缘全厚度控制在480℃～580℃的温度范围内，在空冷阶段充分析出细小的针状铁素体。
如翼缘表面冷却速度低于5℃/s，则芯部冷却速度降低，将析出呈条带状分布的块状铁素体或珠光体，损害韧性和塑性，设定下限为5℃/s；如冷却速度高于13℃/s，则总冷却时间不足，芯部热传导不充分，空冷起始温度高，将会形成大量珠光体，损害韧性和塑性，设定上限为13℃/s。
如翼缘表面冷却温度低于480℃，则表面及附近区域落入上贝氏体析出区间，将会形成超过3％的贝氏体，而在芯部区域形状针状铁素体，组织差距增大，损害厚度方向性能，设定下限为480℃；如冷却温度高于580℃，则芯部空冷起始温度升高，将会形成析出大量块状铁素，而在表面及附近区域形成宽大的针状铁素体，损害韧性和厚度方向性能，设定上限为580℃。
本发明的超厚规格热轧H型钢，翼缘厚度范围90mm～150mm，此时的腹板厚度范围是50mm～120mm。
进一步可选的，翼缘厚度范围90mm～115mm。
进一步可选的，翼缘厚度范围115mm～150mm。
由于重型支撑结构件的设计需要一定强度和刚度，要求采用的热轧H型钢，其翼缘厚度 不低于90mm，设定下限为90mm；如厚度超过150mm，则需要更大尺寸的异型坯，设备投资大、生产难度高，且翼缘过厚，轧制变形渗透和控制冷却渗透有限，设定上限为150mm。
根据结构设计相关要求和热轧H型钢技术特点，在翼缘厚度处于90mm～150mm范围内时，结构稳定性、生产可实现性方面考虑，腹板厚度为50mm～120mm。
以下表1-表4分别是本发明提供的实施例1-10的化学组分、生产工艺参数、显微组织情况及力学性能情况：
表1本发明实施例1～10的化学成分(单位：wt％)
序号 C Si Mn P S Cr Ni Cu V Nb TI Al N 0 CEV％
1 0.04 0.15 0.81 0.010 0.004 0.71 0.03 0.11 0.047 0.044 0.011 0.014 0.0075 0.0042 0.34
0.05 0.19 0.85 0.013 0.003 0.80 0.01 0.13 0.042 0.047 0.013 0.012 0.0070 0.0039 0.37
3 0.07 0.26 0.90 0.014 0.004 0.84 0.01 0.17 0.051 0.053 0.015 0.022 0.0089 0.0041 0.41
4 0.06 0.21 0.95 0.011 0.003 0.89 0.02 0.20 0.048 0.055 0.014 0.020 0.0083 0.0055 0.42
0.08 0.30 0.62 0.011 0.003 0.92 0.02 0.27 0.086 0.029 0.022 0.020 0.0110 0.0021 0.40
6 0.09 0.33 0.69 0.013 0.003 0.95 0.01 0.24 0.085 0.033 0.021 0.025 0.0113 0.0023 0.43
7 0.10 0.31 0.75 0.012 0.003 0.93 0.03 0.33 0.089 0.039 0.023 0.027 0.0120 0.0036 0.45
8 0.11 0.39 0.73 0.014 0.002 0.96 0.01 0.35 0.089 0.038 0.025 0.023 0.0119 0.0031 0.47
9 0.10 0.35 0.79 0.010 0.003 0.95 0.03 0.39 0.096 0.040 0.023 0.029 0.0120 0.0029 0.47
10 0.11 0.38 0.72 0.011 0.002 0.97 0.02 0.39 0.093 0.040 0.024 0.024 0.0118 0.0030 0.47
表2本发明实施例1～10的主要工艺参数
Figure PCTCN2021126546-appb-000001
表3本发明实施例1～10的显微组织情况
Figure PCTCN2021126546-appb-000002
本发明的超厚规格热轧H型钢，根据标准GB/T 2975规定，在翼缘上，以距离端部的宽 度1/6处及厚度1/4处取样；根据标准GB/T 228.1规定，测得的室温拉伸屈服强度应不低于460MPa、抗拉强度应不低于540MPa，断后伸长率应不低于24.0％；根据标准GB/T 229规定，测得的-20℃冲击功值应不低于80J；根据标准GB/T 5313规定，测得的厚度方向性能应达到Z35级别。
表4本发明实施例1～10的力学性能情况
Figure PCTCN2021126546-appb-000003
由表1～表4提供的实施例可知，采用本发明所述方法生产翼缘厚度90mm～150mm的超厚规格热轧H型钢，室温屈服强度达到464MPa～522MPa，抗拉强度597MPa～649MPa，断后伸长率24.0％～32.0％，-20℃冲击功为84J～126J，厚度方向性能超过Z35级别要求。
以上仅为本发明的较佳实施例，但本发明的保护范围并不局限于此，任何熟悉本技术领域的技术人员在本发明揭露的技术范围内，可轻易想到的变化或替换，都应涵盖在本发明的保护范围之内。因此，本发明的保护范围应该以权利要求所界定的保护范围为准。
本发明未详述之处，均为本技术领域技术人员的公知技术。

Claims (10)

  1. 一种超厚规格热轧H型钢，其特征在于：所述H型钢的化学成分按质量百分比计，包括C：0.04～0.11，Si：0.10～0.40，Mn：0.40～1.00，Cr：0.40～1.00，Cu：0.10～0.40，Nb：0.020～0.060，V：0.040～0.100，Ti：0.010～0.025，Al：0.010～0.030，N：0.0060～0.0120，P：≤0.015，S：≤0.005，O：≤0.0060，满足0.090％≤Nb+V+Ti≤0.170％和6.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.30％≤CEV≤0.48％。
  2. 根据权利要求1所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的化学成分按质量百分比计，包括C：0.04～0.07，Si：0.10～0.30，Mn：0.80～1.00，Cr：0.40～0.90，Cu：0.10～0.25，Nb：0.040～0.060，V：0.040～0.080，Ti：0.010～0.015，Al：0.010～0.020，N：0.0060～0.0100，P：≤0.015，S：≤0.005，O：≤0.0060，满足0.090％≤Nb+V+Ti≤0.130％，6.5≤(V+Ti)/N≤8.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.30％≤CEV≤0.43％。
  3. 根据权利要求1所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的化学成分按质量百分比计，包括C：0.07～0.11，Si：0.30～0.40，Mn：0.40～0.80，Cr：0.90～1.00，Cu：0.25～0.40，Nb：0.020～0.040，V：0.080～0.100，Ti：0.015～0.025，Al：0.020～0.030，N：0.0100～0.0120，P：≤0.015，S：≤0.005，O：≤0.0040，满足0.130％＜Nb+V+Ti≤0.170％，8.5≤(V+Ti)/N≤10.5，其余为Fe及微量残余元素，根据CEV＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu)/15计算，化学成分满足0.40％≤CEV≤0.48％。
  4. 根据权利要求1至3任意一项所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，显微组织以面积百分比计，包含85％～98％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于2％，其中铁素体晶粒宽度尺寸不大于40μm，沿翼缘厚度方向不同区域的针状铁素体含量差距不大于16％。
  5. 根据权利要求1至3任意一项所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，显微组织以面积百分比计，包含85％～91％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于2％，其中铁素体晶粒宽度尺寸不大于20μm，沿翼缘厚度方向不同区域的针状铁素体含量差距不大于9％。
  6. 根据权利要求1至3任意一项所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，显微组织以面积百分比计，包含91％～98％的针状铁素体，剩余组织为贝氏体或残余奥氏体，贝氏体含量不高于1％，其中铁素体晶粒宽度尺寸20μm～40μm，沿翼缘厚度方向不同区域的针状铁素体含量差距9％～16％。
  7. 根据权利要求1至3任意一项所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的翼缘距离端部的宽度1/6处、厚度1/4处，室温拉伸屈服强度不低于460MPa，抗拉强度不低于540MPa，断后伸长率不低于24.0％；-20℃冲击功值不低于80J，厚度方向性能达到Z35级别。
  8. 根据权利要求1至3任意一项所述的一种超厚规格热轧H型钢，其特征在于：所述H型钢的翼缘厚度为90mm～150mm。
  9. 一种如权利要求1-8任意一项所述的超厚规格热轧H型钢的生产方法，其特征在于，包括以下步骤：坯料加热，加热温度1200℃～1350℃，加热时间120min～180min；开坯轧制，轧制完成后翼缘表面温度不低于1000℃；以不低于20℃/s的冷却速度喷水冷却，将翼缘表面快速冷却至700℃～800℃，然后进入万能轧机轧制，万能轧制完成后，先采用喷水冷却的方式，以5℃/s～13℃/s的冷却速度，将轧件翼缘表面快速冷却至480℃～530℃，然后空冷。
  10. 根据权利要求9所述的超厚规格热轧H型钢的生产方法，其特征在于：所述的异型坯开坯轧制，轧制完成后翼缘表面温度不低于1020℃。
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