WO2015143932A1 - 一种屈服强度890MPa级低焊接裂纹敏感性钢板及其制造方法 - Google Patents

一种屈服强度890MPa级低焊接裂纹敏感性钢板及其制造方法 Download PDF

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WO2015143932A1
WO2015143932A1 PCT/CN2015/070729 CN2015070729W WO2015143932A1 WO 2015143932 A1 WO2015143932 A1 WO 2015143932A1 CN 2015070729 W CN2015070729 W CN 2015070729W WO 2015143932 A1 WO2015143932 A1 WO 2015143932A1
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steel plate
rolling
temperature
steel
yield strength
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PCT/CN2015/070729
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English (en)
French (fr)
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姚连登
赵四新
王鹏建
苗雨川
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宝山钢铁股份有限公司
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Priority to KR1020167026018A priority Critical patent/KR102291866B1/ko
Priority to BR112016021752-7A priority patent/BR112016021752B1/pt
Priority to EP15767692.5A priority patent/EP3124640B1/en
Priority to US15/128,970 priority patent/US20180355452A1/en
Priority to JP2016558640A priority patent/JP6502377B2/ja
Priority to AU2015235813A priority patent/AU2015235813A1/en
Publication of WO2015143932A1 publication Critical patent/WO2015143932A1/zh

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0231—Warm rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese

Definitions

  • the present invention relates to a high strength low weld crack sensitive steel sheet, and more particularly to a low weld crack sensitive steel sheet having a yield strength of 890 MPa and a method of manufacturing the same.
  • thermomechanical treatment of steel sheets is usually carried out by controlled rolling and cooling (TMCP). By controlling the deformation rate and the cooling rate, the microstructure is refined or a high-strength structure such as ultra-fine bainite is formed to increase the yield strength of the steel.
  • the components of low-carbon high-strength steel produced by TMCP are mainly Mn-Ni-Nb-Mo-Ti and Si-Mn-Cr-Mo-Ni-Cu-Nb-Ti-Al-B systems.
  • a low-alloy high-strength steel produced by the TMCP process at two temperature stages published by International Publication No. WO99/05335, has a chemical composition (wt.%) of C: 0.05 to 0.10% and Mn: 1.7 to 2.1%.
  • Another example is an ultra-low carbon bainite steel disclosed in Chinese Patent Publication No. 1521285, whose chemical composition (wt.%) is C: 0.01-0.05%, Si: 0.05-0.55%, Mn: 1.0-2.2%, Ni: 0.0 to 1.0%, Mo: 0.0 to 0.5%, Cr: 0.0 to 0.7%, Cu: 0.0 to 1.8%, Nb: 0.015 to 0.070%, Ti: 0.005 to 0.03%, B: 0.0005 to 0.005%, Al : 0.015 to 0.07%.
  • the alloy elements of the two steel types disclosed above are designed as Mn-Ni-Nb-Mo-Ti and Si-Mn-Cr-Mo-Ni-Cu-Nb-Ti-Al-B, respectively, since both Mo and Ni are Valuable alloys, therefore, from the type of alloying elements added and the total amount added, the cost of preparing such steels is relatively high.
  • the object of the present invention is to provide a low-weld crack sensitive steel plate with a yield strength of 890 MPa and a manufacturing method thereof, which adopts Si-Mn-Nb-Mo-V-Ti-Al-B steel grade and controls thermomechanical rolling. And cooling technology, and no quenching and tempering treatment, its welding crack sensitivity index Pcm ⁇ 0.25%, yield strength are greater than 890MPa, tensile strength is greater than 950MPa, Xia's impact energy Akv (-20 ° C) ⁇ 120J, plate thickness 60mm, with good low temperature toughness and weldability, low-carbon ultra-fine bainitic lath low weld crack sensitive steel plate.
  • a low-weld crack-sensitive steel plate with a yield strength of 890 MPa the chemical composition weight percentage thereof being: C 0.06 to 0.13 wt.%, Si 0.05 to 0.70 wt.%, Mn 1.20 to 2.30 wt.%, Mo0 to 0.25 wt.% , Nb 0.03 to 0.11 wt.%, Ti 0.002 to 0.050 wt.%, Al 0.02 to 0.15 wt.%, B0 to 0.0020 wt.%, 2Si+3Mn+4Mo ⁇ 8.5, and the balance being Fe and unavoidable impurities;
  • the steel plate satisfies the welding crack sensitivity index Pcm ⁇ 0.25%.
  • composition design of the present invention is a composition design of the present invention:
  • C enlarges the austenite region, and carbon in the supersaturated ferrite structure formed by quenching increases its strength. C is detrimental to the welding performance. The higher the C content, the worse the welding performance. For the bainitic steel produced by the TMCP process, the lower the C content, the better the toughness, and the lower the carbon content, the higher the toughness steel plate can be produced. Therefore, the present invention
  • the C content is controlled to be 0.06 to 0.13%.
  • Si:Si does not form carbides in steel, but exists in solid solution form in bainitic ferrite or austenite. It increases the strength of bainite austenite or ferrite in steel.
  • the solid solution strengthening effect of Si is stronger than that of Mn, Nb, Cr, W, Mo and V.
  • Si reduces the diffusion rate of carbon in the austenite, and shifts the CCT curve ferrite and pearlite C curves to the right, which is favorable for the formation of bainite structure during continuous cooling.
  • the addition of not more than 0.70% of Si to the steel of the present invention is advantageous for improving the strength and toughness matching relationship of the steel.
  • Mn and Fe form a solid solution, which improves the strength and hardness of bainitic ferrite and austenite in steel.
  • Mn enlarges the austenite region in the iron-carbon equilibrium phase diagram, which makes the steel's ability to form a stable austenite structure second only to Ni, which strongly increases the hardenability of the steel.
  • the Mn content is high, there is a tendency that the steel grains are coarsened.
  • 1.20 to 2.30% of Mn is added to slow the transformation of ferrite and pearlite, which is advantageous for forming a refined bainite structure and imparting a certain strength to the steel.
  • Mo and Cr are ferrite elements, which shrink the austenite region. Mo and Cr are solid-solved in austenite and ferrite to increase their strength, improve the hardenability of steel, and prevent temper brittleness. Mo is a very expensive element, and the present invention does not require tempering and quenching treatment. The present invention only needs to add not more than 0.25% of Mo and no more than 0.20 of Cr to achieve cost reduction.
  • Nb The present invention achieves the purpose of refining crystal grains and increasing the thickness of the steel sheet by adding more Nb, and on the other hand, increasing the non-recrystallization temperature of the steel, and facilitating the relative use in the rolling process. Higher finishing temperature, which speeds up the rolling speed and increases production efficiency. In addition, the thickness of the steel sheet that can be produced is increased due to enhanced grain refining. In the present invention, 0.03 to 0.10 wt.% of Nb is added, taking into account the solid solution strengthening and fine grain strengthening of Nb.
  • Ti is a ferrite element that strongly reduces the austenite region.
  • the Ti carbide of Ti is relatively stable and can suppress grain growth.
  • Ti is solid-dissolved in austenite, which contributes to improved hardenability of steel.
  • Ti can reduce the temper brittleness of the first type of 250 to 400 ° C, but the present invention does not require quenching and tempering treatment, so the amount of Ti added can be reduced. In the present invention, 0 to 0.050 wt.% is added to form a fine carbonitride precipitate, and the bainite lath is refined.
  • Al increases the phase change driving force of austenite to ferrite transformation, and is an element that strongly reduces the austenite phase circle. Al interacts with N in steel to form fine and diffused AlN precipitates, which can inhibit grain growth, refine grains and improve the toughness of steel at low temperatures. Excessive Al content has an adverse effect on the hardenability and weldability of the steel. In the present invention, not more than 0.15% of Al refined grains are added to improve the toughness and ensure the welding performance of the steel sheet.
  • B can significantly increase the hardenability of steel.
  • the addition of 0 to 0.002% of B in the present invention makes it possible to obtain a high-strength bainite structure relatively easily under certain cooling conditions.
  • the content of the three elements of Si, Mn and Mo should be in accordance with the following relationship: 2Si+3Mn+4Mo ⁇ 8.5, to satisfy the good welding performance of the steel sheet of the invention. Specifically, it is possible to ensure that the steel plate having a thickness of 60 mm or less is welded without cracks at a relatively low preheating temperature (normal temperature to 50 ° C).
  • the action of various alloying elements can be reasonably utilized to produce a steel plate having a maximum thickness of 60 mm.
  • the weld crack sensitivity index Pcm of a low weld crack sensitive steel sheet can be determined by the following formula:
  • the welding crack sensitivity index Pcm is a judgment index reflecting the tendency of the steel to weld cold cracks, and the lower the Pcm, the better the weldability, and conversely, the weldability is worse.
  • Good weldability means that weld cracks are not easily generated during welding, and steel with poor weldability is prone to cracks.
  • steel In order to avoid cracks, steel must be preheated before welding. The better the weldability, the more preheating temperature is required. Low, otherwise a higher preheat temperature is required.
  • the steel grade of Q800CF shall have a Pcm value of less than 0.28%.
  • the high-strength low-weld crack-sensitive steel sheet of the ultra-fine bainite strip according to the present invention has a weld crack sensitivity of less than 0.20% and has excellent weldability.
  • a method for manufacturing a low-weld crack-sensitive steel sheet having a yield strength of 890 MPa Including the following steps:
  • the slab or ingot is smelted and cast according to the following composition, the thickness of which is not less than 4 times the thickness of the finished steel plate; the chemical composition weight percentage is: C 0.06-0.13wt.%, Si 0.05-0.70wt.%, Mn 1.20 ⁇ 2.30wt.%, Mo 0-0.25wt.%, Nb0.03 ⁇ 0.11wt.%, Ti 0.002 ⁇ 0.050wt.%, Al 0.02 ⁇ 0.15wt.%, B0 ⁇ 0.0020wt.%, 2Si+3Mn+ 4Mo ⁇ 8.5, the rest is Fe and unavoidable impurities; and, the steel plate satisfies the welding crack sensitivity index Pcm ⁇ 0.25%;
  • Heating temperature is 1050 ⁇ 1180 ° C, holding time is 120 ⁇ 180 minutes;
  • Rolling is divided into first stage and second stage rolling;
  • the rolling temperature is 1050 ⁇ 1150 ° C, when the thickness of the rolled piece reaches 2 to 3 times the thickness of the finished steel plate, the temperature on the roller table is to be 800 ⁇ 860 ° C;
  • the pass deformation rate is 10 to 28%, and the finishing temperature is 780 to 840 ° C;
  • the steel plate is cooled to 220 to 350 ° C at a rate of 15 to 30 ° C / S, and air-cooled after effluent.
  • step 3 the air cooling is cooled by stacking or cooling.
  • the non-recrystallization temperature is about 950 to 1050 °C.
  • the temperature of the rolled billet is reduced to 800-860 °C, the recovery and static recrystallization process occur in the austenite grains, and the austenite grains are refined.
  • carbonitrides of Nb, V and Ti are separately precipitated and composited.
  • the precipitated carbonitrides pinned the dislocation and subgrain boundary motion, retained a large number of dislocations in the austenite grains, and provided a large number of nucleation sites for the formation of bainite during cooling. Rolling at 800 to 860 ° C greatly increases the dislocation density in austenite.
  • the carbonitride precipitated on the dislocation suppresses the coarsening of the crystal grains after the deformation. Due to the effect of deformation-induced precipitation, a larger ductal deformation rate will favor the formation of finer and dispersed precipitates.
  • High-density dislocations and finely dispersed precipitates provide a high-density nucleation site for bainite, and pinning of the bainite-grown interface by the second-phase particles inhibits the growth and coarsening of bainite laths. This pair Both the strength and toughness of the steel play an advantageous role.
  • the finishing temperature is controlled in the low temperature section of the non-recrystallization zone, and the temperature zone is close to the phase transition point Ar 3 , that is, the finishing rolling temperature is 780-840 ° C, and the final rolling in this temperature range can increase the deformation and suppress the recovery. It increases the defects in austenite, provides higher energy accumulation for bainite transformation, and does not bring excessive load to the rolling mill, which is more suitable for thick plate production.
  • the steel sheet After the end of the rolling, the steel sheet enters the accelerated cooling device and is cooled to 450 to 550 ° C at a rate of 15 to 30 ° C / sec.
  • the faster cooling rate avoids the formation of ferrite and pearlite and directly enters the bainite transformation zone of the CCT curve.
  • the bainite phase change driving force can be expressed as:
  • ⁇ G chem is the chemical driving force
  • ⁇ G d is the strain storage energy caused by the defect.
  • the large cooling rate makes the austenite supercooled, increasing the driving force of the chemical phase change, and considering the strain storage energy ⁇ G d caused by the rolling process, the driving force of the bainite nucleation is increased. Due to the high dislocation density in the grains, the nucleation sites of bainite increase. In combination with thermodynamics and kinetics, bainite nucleates at a large rate. The faster cooling rate allows the bainite transformation to be completed very quickly, inhibiting the coarsening of the bainitic ferrite lath.
  • the air-cooling of the stack can make the precipitation of V in the ferrite more complete, and increase the contribution of precipitation strengthening to the strength. Therefore, by using the heat treatment method of the present invention, a matrix structure mainly composed of refined bainite can be obtained, and a steel sheet having high strength and good toughness can be produced.
  • thermomechanical treatment of steel sheets is usually carried out by controlled rolling and cooling (TMCP).
  • TMCP controlled rolling and cooling
  • the microstructure is refined or a high-strength structure such as ultra-fine bainite is formed to increase the yield strength of the steel.
  • the microalloying element Nb is added to the component of the present invention, and Nb can form a carbonitride during the heat treatment, and has a precipitation strengthening effect. Nb dissolved in the matrix has a solid solution strengthening effect.
  • TMCP and relaxation controlled precipitation (RPC) techniques are used in the heat treatment to form a stable dislocation network, which disperses fine second phase particles at dislocations and subgrain boundaries, and promotes nucleation and inhibition.
  • the growth of the bainite lath is realized, and the combination of dislocation strengthening, precipitation strengthening and fine grain strengthening is formed, which improves the strength and toughness of the steel.
  • the basic principle is as follows:
  • the steel sheet is sufficiently deformed in the recrystallization zone to cause high defect accumulation in the deformed austenite, which greatly increases the dislocation density in the austenite.
  • the recovery and recrystallization that occur during the rolling process refine the prior austenite grains.
  • the intragranular dislocations are rearranged. Due to the presence of a hydrostatic pressure field in the edge dislocations, interstitial atoms such as B are enriched toward dislocations, grain boundaries and subgrain boundaries, reducing dislocation mobility.
  • the high-density dislocations caused by the deformation evolved during the recovery process forming a stable dislocation network.
  • microalloying elements such as Nb, V, and Ti are precipitated at the grain boundary, subgrain boundary, and dislocation with carbonitrides of different stoichiometric ratios such as (Nb, V, Ti) x (C, N) y .
  • Two-phase particles such as precipitated carbonitrides pin the dislocations and subgrain boundaries in the crystal grains, stabilizing substructures such as dislocation walls.
  • the rolling after relaxation further increases the dislocation density in the steel.
  • the deformed austenite is relaxed, the deformed austenite grains with the dislocation and carbonitride precipitation configuration are in the phase transition at the beginning, and are not relaxed after deformation and a large number of dislocations are scattered. different.
  • the subgrain boundary with a certain orientation difference is a nucleation preferential position, and if a second phase having a different phase interface with the matrix exists in the vicinity thereof, it is more advantageous for the new phase nucleation at the time of phase change.
  • a large number of new phase grains will nucleate within the original austenite grains.
  • the orientation difference between the subcrystals is increased to some extent.
  • the intermediate temperature transition products, such as bainite are hindered by the front subgrain boundary during the growth process after nucleation at the subgrain boundary.
  • bainitic ferrite When bainitic ferrite is formed, its phase transition interface is dragged by the precipitated second phase carbonitride particles, which inhibits the growth process.
  • the TMCP+RPC process forms a high-density dislocation network structure and the second phase precipitates a large number of potential nucleation sites for the nucleation of bainitic ferrite.
  • the drag effect of the second phase particles on the moving interface and the evolution of the subgrain boundary on the growth of the bainite The combined effect of the process on promoting nucleation and inhibiting growth refines the bainitic ferrite slabs of the final structure.
  • the C content is greatly reduced, and part of Mo is replaced by an inexpensive alloying element such as Mn, and the fine precipitated particles of C and N of Nb are used for precipitation strengthening instead of Cu.
  • Precipitation strengthening does not require the addition of precious elements such as Ni, and the alloying element content is low, the raw material cost is low, the welding crack sensitivity is small, and no preheating is required before welding.
  • the steel sheet of the present invention does not require any additional heat treatment, thereby simplifying the manufacturing process and reducing the manufacturing cost of the steel.
  • the low weld crack sensitive steel plate of the present invention has a yield strength greater than 890 MPa, a tensile strength greater than 950 MPa, a Charpy impact energy Akv (-20 ° C) ⁇ 100 J, and a plate thickness of up to 60 mm.
  • the welding crack sensitivity index Pcm ⁇ 0.25% has excellent welding performance.
  • the present invention can prepare thick plates having a maximum thickness of 60 mm.
  • Table 1 shows the chemical composition (wt.%) and the Pcm (%) value of the steel sheet according to the examples of the present invention.
  • Table 2 shows the mechanical properties of the steel sheets of the examples of the present invention.
  • Table 3 shows the results of the welding performance test (small iron test) of the 890 MPa class low weld crack sensitive steel sheet according to Example 1 of the present invention.
  • the embodiment is the same as the first embodiment, wherein the heating temperature is 1050 ° C, and the temperature is maintained for 240 minutes; the first stage rolling has an opening rolling temperature of 1040 ° C, the rolling piece thickness is 90 mm; and the second stage rolling has an opening rolling temperature of 840 ° C.
  • the pass rate is 15-20%, the finish rolling temperature is 810 ° C, the finished steel plate thickness is 30 mm, the steel plate cooling rate is 25 ° C / S, and the termination temperature is 350 ° C.
  • the embodiment is the same as the first embodiment, wherein the heating temperature is 1150 ° C, and the temperature is maintained for 150 minutes; the first stage rolling has an opening rolling temperature of 1080 ° C, the rolling piece thickness is 120 mm; and the second stage rolling has an opening rolling temperature of 830 ° C.
  • the pass deformation rate is 10 to 15%, the finish rolling temperature is 820 ° C, the finished steel plate thickness is 40 mm, the steel plate cooling rate is 20 ° C / S, and the termination temperature is 330 ° C.
  • the embodiment is the same as the first embodiment, wherein the heating temperature is 1120 ° C, and the temperature is maintained for 180 minutes; the first stage rolling has an opening rolling temperature of 1070 ° C, the rolling piece thickness is 150 mm; and the second stage rolling has an opening rolling temperature of 830 ° C.
  • the pass deformation rate is 10 to 20%, the finish rolling temperature is 800 ° C, the finished steel plate thickness is 50 mm, the steel plate cooling rate is 15 ° C / S, and the termination temperature is 285 ° C.
  • the embodiment is the same as the first embodiment, wherein the heating temperature is 1130 ° C, and the temperature is maintained for 180 minutes; the first stage rolling has an opening rolling temperature of 1080 ° C, the rolling piece thickness is 150 mm; and the second stage rolling has an opening rolling temperature of 840 ° C.
  • the pass deformation rate is 10 to 15%, the finish rolling temperature is 810 ° C, the finished steel plate thickness is 60 mm, the steel plate cooling rate is 15 ° C / S, and the termination temperature is 220 ° C.
  • the embodiment is the same as the first embodiment, wherein the heating temperature is 1120 ° C, and the temperature is maintained for 180 minutes; the first stage rolling has an opening rolling temperature of 1050 ° C, the rolling piece thickness is 120 mm; and the second stage rolling has an opening rolling temperature of 820 ° C.
  • the pass deformation rate is 15 to 25%, the finish rolling temperature is 780 ° C, the finished steel plate thickness is 40 mm, the steel plate cooling rate is 20 ° C / S, and the termination temperature is 300 ° C.
  • Example C Si Mn Nb Al Ti Cr Mo B Fe Pcm 1 0.09 0.35 1.80 0.070 0.02 0.015 0.16 0.25 0.0018 the remaining 0.217 2 0.06 0.70 2.25 0.045 0.06 0.020 0 0 0.0010 the remaining 0.201 3 0.08 0.40 2.06 0.085 0.04 0.050 0.20 0.10 0.0011 the remaining 0.218
  • the present invention relates to a low-weld crack-sensitive steel plate with a yield strength of 890 MPa, Pcm ⁇ 0.25%, a yield strength of more than 890 MPa, a tensile strength of more than 950 MPa, and a Charpy impact force Akv (-20).
  • °C) ⁇ 120J plate thickness up to 60mm, with good low temperature toughness and weldability.
  • the welding performance test (small iron test) of the steel sheet of Example 1 of the present invention was carried out, and no crack was found at room temperature and 50 ° C (see Table 3), indicating that the steel of the present invention has good welding performance and is generally not welded. Need to warm up.

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Abstract

一种屈服强度890MPa级低焊接裂纹敏感性钢板及其制造方法，其成分重量百分比为：C：0.06～0.13wt.％、Si：0.05～0.70wt.％、Mn：1.20～2.30wt.％、Mo：0～0.25wt.％、Nb：0.03～0.11wt.％、Ti：0.002～0.050wt.％、Al：0.02～0.15wt.％、B：0～0.0020wt.％，2Si+3Mn+4Mo≤8.5，其余为Fe和不可避免的杂质。采用控制热机械轧制和冷却技术，有利于钢板强度﹑塑性和韧性的提高。该钢板的屈服强度大于890MPa、抗拉强度大于950MPa、夏氏冲击功Akv(-20℃)≥120J，焊接裂纹敏感性指数Pcm≤0.25％。

Description

一种屈服强度890MPa级低焊接裂纹敏感性钢板及其制造方法 技术领域
本发明涉及高强度低焊接裂纹敏感性钢板，具体地说，本发明涉及一种屈服强度890MPa级低焊接裂纹敏感性钢板及其制造方法。
背景技术
高强度机械设备及工程建设用钢，需要较高的强度和优良的韧性，各种因素对强度的贡献可以用下式表示：
σ＝σf+σp+σsl+σd
式中σf是细晶强化，σp是析出强化，σsl是固溶强化，σd是位错强化。钢板的热机械处理通常采用控轧控冷方式(TMCP)。通过控制变形率和冷却速度实现微观组织的细化或形成超细贝氏体等高强度组织，提高钢的屈服强度。
目前，采用TMCP生产低碳高强钢的成分主要是Mn-Ni-Nb-Mo-Ti和Si-Mn-Cr-Mo-Ni-Cu-Nb-Ti-Al-B体系。
如国际公布号为WO99/05335公布的一种以两个温度阶段以TMCP工艺生产的低合金高强钢，其化学成分(wt.％)为：C：0.05～0.10％、Mn：1.7～2.1％、Ni：0.2～1.0％、Mo：0.25～0.6Mo％、Nb：0.01～0.10％、Ti：0.005～0.03％、P≤0.015％、S≤0.003％。
又如中国专利公开号为1521285公布的一种超低碳贝氏体钢，其化学成分(wt.％)为C：0.01～0.05％、Si：0.05～0.55％、Mn：1.0～2.2％、Ni：0.0～1.0％、Mo：0.0～0.5％、Cr：0.0～0.7％、Cu：0.0～1.8％、Nb：0.015～0.070％、Ti：0.005～0.03％、B：0.0005～0.005％、Al：0.015～0.07％。
上述公开的两种钢种的合金元素设计分别为Mn-Ni-Nb-Mo-Ti和Si-Mn-Cr-Mo-Ni-Cu-Nb-Ti-Al-B体系，由于Mo和Ni均为贵重合金，因此从添加的合金元素的种类和加入的总量来分析，制备此类钢种成本较高。
发明内容
本发明的目的在于提供一种屈服强度890MPa级低焊接裂纹敏感性钢板及其制造方法，采用Si-Mn-Nb-Mo-V-Ti-Al-B系钢种，通过控制热机械轧制 和冷却技术，且无调质处理，其焊接裂纹敏感性指数Pcm≤0.25％，屈服强度均大于890MPa，抗拉强度大于950MPa，夏氏冲击功Akv(－20℃)≥120J，板厚可达60mm，具有良好的低温韧性和焊接性，为低碳超细贝氏体板条低焊接裂纹敏感性钢板。
为达到上述目的，本发明的技术方案是：
一种屈服强度890MPa级低焊接裂纹敏感性钢板，其化学成分重量百分比为：C 0.06～0.13wt.％、Si 0.05～0.70wt.％、Mn 1.20～2.30wt.％、Mo0～0.25wt.％、Nb 0.03～0.11wt.％、Ti 0.002～0.050wt.％、Al 0.02～0.15wt.％、B0～0.0020wt.％，2Si+3Mn+4Mo≤8.5，其余为Fe和不可避免的杂质；且，钢板满足焊接裂纹敏感性指数Pcm≤0.25％。
在本发明的成分设计中：
C：C扩大奥氏体区，淬火形成的过饱和铁素体组织中的碳可增加其强度。C对焊接性能不利。C含量越高，焊接性能越差，对于采用TMCP工艺生产的贝氏体钢来说，C含量越低则韧性越好，较低的碳含量可以生产更大厚度的高韧性钢板，因此本发明C含量控制为0.06～0.13％。
Si：Si在钢中不形成碳化物，而是以固溶形态存在于贝氏体铁素体或奥氏体中。它提高钢中贝氏体奥氏体或铁素体的强度。Si的固溶强化作用较Mn、Nb、Cr、W、Mo和V强。Si降低奥氏体中碳的扩散速度，使CCT曲线铁素体和珠光体C曲线向右移动，有利于连续冷却过程中形成贝氏体组织。本发明钢中加入不超过0.70％的Si有利于提高钢的强度和韧性匹配关系。
Mn：Mn和Fe可形成固溶体，提高钢中贝氏体铁素体和奥氏体的强度和硬度。Mn扩大铁碳平衡相图中的奥氏体区，它使钢形成稳定奥氏体组织的能力仅次于Ni，强烈增加钢的淬透性。Mn含量较高时，有使钢晶粒粗化的倾向。本发明中加入1.20～2.30％的Mn，减缓铁素体和珠光体转变的速度，有利于形成细化的贝氏体组织，并使钢具有一定的强度。
Mo和Cr：Mo和Cr是铁素体化元素，缩小奥氏体区。Mo、Cr固溶在奥氏体和铁素体中提高其强度，提高钢的淬透性，防止回火脆性。Mo是一种十分昂贵的元素，且本发明无需回火调质处理，本发明只需加入不超过0.25％的Mo和不超过0.20的Cr，以达到降低成本的目的。
Nb：本发明通过加入较多的Nb，一方面以达到细化晶粒和增加钢板厚度的目的，另一方面是提高钢的未再结晶温度，便于在轧制过程中采用相对 较高的终轧温度，从而加快轧制速度，提高生产效率。此外，由于强化了晶粒细化作用，使得可生产钢板的厚度增大。本发明中加入0.03～0.10wt.％的Nb，兼顾了Nb的固溶强化和细晶强化作用。
Ti：Ti是铁素体化元素，强烈缩小奥氏体区。Ti的碳化物TiC比较稳定，可以抑制晶粒长大。Ti固溶在奥氏体中，有利于提高的钢的淬透性。Ti可降低第一类250～400℃回火脆性，但本发明不需要调质处理，所以可减少了Ti的添加量。本发明中加入0～0.050wt.％，形成细微的碳氮化物析出，细化贝氏体板条。
Al：Al可增加奥氏体向铁素体转变的相变驱动力，是强烈缩小奥氏体相圈的元素。Al在钢中与N相互作用，形成细小而弥散的AlN析出，可抑制晶粒长大，达到细化晶粒、提高钢在低温下的韧性的目的。Al含量过大对钢的淬透性和焊接性能有不利影响。本发明中加入不超过0.15％的Al细化晶粒，提高韧性并保证钢板的焊接性能。
B：B能够显著增加钢的淬透性，本发明加入0～0.002％的B，可以使钢在一定冷却条件下，比较容易地获得高强度贝氏体组织。
Si、Mn、Mo三种元素含量之间应符合以下关系式：2Si+3Mn+4Mo≤8.5，以满足本发明钢板具有良好的焊接性能。具体来讲，可以保证60mm及以下厚度钢板在较低预热温度(常温至50℃)条件下，焊接无裂纹。
采用本发明设计的化学成分，可合理利用各种合金元素的作用，生产最大厚度为60mm的钢板。
低焊接裂纹敏感性钢板的焊接裂纹敏感性指数Pcm可按下式确定：
Pcm＝C+Si/30+Ni/60+(Mn+Cr+Cu)/20+Mo/15+V/10+5B
焊接裂纹敏感性指数Pcm是反映钢的焊接冷裂纹倾向的判定指标，Pcm越低，焊接性越好，反之，则焊接性越差。焊接性好是指焊接时不易产生焊接裂纹，而焊接性差的钢容易产生裂纹，为了避免裂纹的产生，必须在焊接前对钢进行预热，焊接性越好，则所需的预热温度越低，反之则需要较高的预热温度。根据中华人民共和国黑色冶金行业标准YB/T 4137－2005规定，牌号为Q800CF的钢种，Pcm值需低于0.28％。本发明所涉及超细贝氏体板条高强度低焊接裂纹敏感性钢板的焊接裂纹敏感性低于0.20％，具有优良的焊接性能。
本发明的一种屈服强度890MPa级低焊接裂纹敏感性钢板的制造方法， 包括如下步骤：
1)冶炼、浇铸
按下述成分冶炼、浇铸连铸坯或钢锭，其厚度不小于成品钢板厚度的4倍；其化学成分重量百分比为：C 0.06～0.13wt.％、Si0.05～0.70wt.％、Mn 1.20～2.30wt.％、Mo 0～0.25wt.％、Nb0.03～0.11wt.％、Ti 0.002～0.050wt.％、Al 0.02～0.15wt.％、B0～0.0020wt.％，2Si+3Mn+4Mo≤8.5，其余为Fe和不可避免的杂质；且，钢板满足焊接裂纹敏感性指数Pcm≤0.25％；
2)加热、轧制
加热温度为1050～1180℃，保温时间为120～180分钟；
轧制分为第一阶段和第二阶段轧制；
在第一阶段轧制过程中，开轧温度为1050～1150℃，当轧件厚度到达成品钢板厚度的2～3倍时，在辊道上待温至800～860℃；
在所述第二阶段轧制过程中，道次变形率为10～28％，终轧温度为780～840℃；
3)冷却
钢板以15～30℃/S的速度冷却至220～350℃，出水后空冷。
进一步，步骤3)中，空冷采用堆垛或冷床冷却。
在本发明制造方法中：
①轧制工艺
轧件厚度到达成品钢板厚度的2～3倍时，在辊道上待温至800～860℃。对于含Nb钢来说，其未再结晶温度约为950～1050℃。先在较高温度轧制，奥氏体中存在了一定的位错密度。将轧制钢坯温度降至800～860℃驰豫过程中，奥氏体晶粒内部发生回复、静态再结晶过程，细化了奥氏体晶粒。驰豫过程中同时有Nb、V和Ti的碳氮化物单独析出和复合析出。析出的碳氮化物钉扎了位错和亚晶界运动，在奥氏体晶粒内保留了大量位错，并为冷却过程中贝氏体的形成提供了大量的形核位置。800～860℃轧制，大大增加了奥氏体中的位错密度。位错上析出的碳氮化物，抑制了变形后晶粒的粗化。由于变形诱导析出的作用，较大的道次变形率将有利于形成更加细小和弥散析出物。高密度的位错和细小弥散的析出物为贝氏体提供高密度的形核位置，第二相粒子对贝氏体长大界面的钉扎作用抑制了贝氏体板条长大和粗化，这对 于钢的强度与韧性都起到有利的作用。
将终轧温度控制在未再结晶区的低温段，同时该温度区接近相变点Ar3，即终轧温度为780～840℃，在这个温度范围内终轧，可通过增加变形、抑制回复，增加奥氏体中的缺陷，为贝氏体相变提供更高的能量累积，也不至于给轧机带来过高的负荷，比较适合于厚板生产。
②冷却工艺
轧制结束后，钢板进入加速冷却装置，按15～30℃/秒的速度冷却至450～550℃。较快的冷却速度可以避免铁素体和珠光体的形成，直接进入CCT曲线的贝氏体转变区。贝氏体相变驱动力可以表示为：
△G＝△Gchem+△Gd
式中△Gchem是化学驱动力，△Gd是缺陷造成的应变储存能。较大冷却速度使奥氏体过冷，增加了化学相变驱动力，结合轧制过程造成的应变储存能△Gd考虑，使贝氏体形核的驱动力增加。由于晶粒中高位错密度，贝氏体的形核位置增加。结合热力学和动力学两个因素考虑，贝氏体以很大的速率形核。较快的冷却速度使贝氏体转变很快完成，抑制了贝氏体铁素体板条的粗化。450～550℃堆跺空冷，可使铁素体中V的碳化物析出更加完全，增加了析出强化对强度的贡献。因此采用本发明的热处理方法可以得到以细化的贝氏体为主的基体组织，生产具有较高的强度和良好的韧性的钢板。
高强度机械设备及工程建设用钢，需要较高的强度和优良的韧性，各种因素对强度的贡献可以用下式表示：
σ＝σf+σp+σsl+σd
式中σf是细晶强化，σp是析出强化，σsl是固溶强化，σd是位错强化。钢板的热机械处理通常采用控轧控冷方式(TMCP)。通过控制变形率和冷却速度实现微观组织的细化或形成超细贝氏体等高强度组织，提高钢的屈服强度。本发明成分中添加了微合金元素Nb，热处理过程中Nb可以形成碳氮化物，有析出强化作用。固溶在基体中的Nb，有固溶强化作用。热处理时采用了改进的TMCP和驰豫控制析出(RPC)技术，形成了稳定的位错网络，在位错和亚晶界处析出了弥散细小的第二相粒子，并通过促进形核和抑制长大实现了贝氏体板条细化，形成了位错强化、析出强化和细晶强化的联合作用，提高了钢的强度和韧性，其基本原理为：
钢板在再结晶区充分变形，使变形奥氏体中产生高的缺陷累计，大幅度提高了奥氏体中的位错密度。轧制过程中发生的回复和再结晶细化了原奥氏体晶粒。轧制变形后控冷驰豫过程中，晶内位错会重新排列。由于刃型位错存在静水压力场，间隙原子如B等会向位错、晶界和亚晶界处富集，降低了位错移动性。变形造成的高密度位错在回复过程中经过演化，形成了稳定的位错网络。驰豫过程中，Nb、V、Ti等微合金元素以(Nb,V,Ti)x(C,N)y等不同化学计量比的碳氮化物在晶界、亚晶界和位错处析出。析出的碳氮化物等二相粒子，钉扎了晶粒中的位错和亚晶界，稳定了如位错墙等亚结构。驰豫后轧制使钢中的位错密度进一步增加。驰豫后变形奥氏体在加速冷却时，具有的位错和碳氮化物析出组态的变形奥氏体晶粒在开始相变时，与变形后不驰豫、大量位错混乱分布的情况不同。首先，有一定取向差的亚晶界是形核优先位置，其附近如果存在与基体有异相界面的第二相析出，则更有利于相变时新相形核。驰豫后大量的新相晶粒将在原奥氏体晶粒内形核。其次，由于驰豫后一定量的位错向亚晶界运动，一定程度上增加了亚晶之间的取向差。中温转变产物如贝氏体在亚晶界形核后，长大过程中受到前方亚晶界的阻碍。贝氏体铁素体形成时，其相变界面受到析出的第二相碳氮化物粒子的拖曳作用，抑制了其长大过程。TMCP+RPC工艺形成高密度位错网络结构和第二相析出质点为贝氏体铁素体的形核提供了大量的潜在形核位置。第二相粒子对运动界面的拖曳作用和经过演化的亚晶界对贝氏体的长大有抑制作用。该工艺对促进形核和抑制长大的联合作用细化了最终组织的贝氏体铁素体板条。
对于机械结构和工程建设所使用的高强钢，需要焊前不预热或稍加预热而不产生裂纹，主要是解决了大型钢结构件的焊接施工问题。降低Pcm的唯一手段就是减少碳和合金元素的加入量，而对于采用淬火+回火工艺生产的高强钢来说，减少碳和合金元素的加入量将不可避免地带来钢强度的降低，采用本发明中改进的TMCP+RPC工艺，则可以弥补这种缺陷。本发明采用的成分体系保证钢板具有高强度和低温韧性，同时焊接裂纹敏感性指数Pcm≤0.20％，具有优良的焊接性能。
本发明的有益效果：
1、通过合理设计化学成分，大幅度降低C含量，并且以Mn等廉价合金元素替代部分Mo，以Nb的C﹑N化合微细析出粒子作沉淀强化，代替Cu的 析出强化作用，无需添加Ni等贵重元素，且合金元素含量少，原料成本较低，焊接裂纹敏感性较小，焊前无需预热。
2、本发明钢板不需进行任何额外的热处理，从而简化了制造工序，降低了钢的制造成本。
3、由于成分和工艺设计合理，从实施效果来看，工艺制度比较宽松，可以在中、厚钢板产线上稳定生产。
4、本发明的低焊接裂纹敏感性钢板屈服强度大于890MPa、抗拉强度大于950MPa、夏氏冲击功Akv(-20℃)≥100J、板厚可达60mm。焊接裂纹敏感性指数Pcm≤0.25％，具有优良的焊接性能。
5、本发明可制备最大厚度达60mm的厚板。
本发明的最佳实施方式
以下用实施例对本发明作更详细的描述。这些实施例仅仅是对本发明最佳实施方式的描述，并不对本发明的范围有任何限制。
表1为本发明实施例钢板的化学成分(wt.％)及其Pcm(％)值。表2为本发明实施例钢板的力学性能。表3为本发明实施例1的890MPa级低焊接裂纹敏感性钢板焊接性能试验(小铁研试验)结果。
实施例1
按表2所示的化学成分电炉或转炉冶炼，并浇铸成连铸坯或钢锭，将连铸坯或钢锭加热至1110℃，保温120分钟，在中、厚轧机上进行第一阶段轧制，开轧温度为1050℃，当轧件厚度为60mm时，在辊道上待温至850℃，随后进行第二阶段轧制，第二阶段轧制道次变形率为15～28％，终轧温度为830℃，成品钢板厚度为20mm。轧制结束后，钢板进入加速冷却(ACC)装置，以30℃/S的速度冷却至300℃，出水后堆垛或冷床冷却。
实施例2
实施方式同实施例1，其中加热温度为1050℃，保温240分钟；第一阶段轧制的开轧温度为1040℃，轧件厚度为90mm；第二阶段轧制的开轧温度为840℃，道次变形率为15～20％，终轧温度为810℃，成品钢板厚度为30mm；钢板冷却速度为25℃/S，终止温度为350℃。
实施例3
实施方式同实施例1，其中加热温度为1150℃，保温150分钟；第一阶段轧制的开轧温度为1080℃，轧件厚度为120mm；第二阶段轧制的开轧温度为830℃，道次变形率为10～15％，终轧温度为820℃，成品钢板厚度为40mm；钢板冷却速度为20℃/S，终止温度为330℃。
实施例4
实施方式同实施例1，其中加热温度为1120℃，保温180分钟；第一阶段轧制的开轧温度为1070℃，轧件厚度为150mm；第二阶段轧制的开轧温度为830℃，道次变形率为10～20％，终轧温度为800℃，成品钢板厚度为50mm；钢板冷却速度为15℃/S，终止温度为285℃。
实施例5
实施方式同实施例1，其中加热温度为1130℃，保温180分钟；第一阶段轧制的开轧温度为1080℃，轧件厚度为150mm；第二阶段轧制的开轧温度为840℃，道次变形率为10～15％，终轧温度为810℃，成品钢板厚度为60mm；钢板冷却速度为15℃/S，终止温度为220℃。
实施例6
实施方式同实施例1，其中加热温度为1120℃，保温180分钟；第一阶段轧制的开轧温度为1050℃，轧件厚度为120mm；第二阶段轧制的开轧温度为820℃，道次变形率为15～25％，终轧温度为780℃，成品钢板厚度为40mm；钢板冷却速度为20℃/S，终止温度为300℃。
表1  单位：重量百分比
实施例 C Si Mn Nb Al Ti Cr Mo B Fe Pcm
1 0.09 0.35 1.80 0.070 0.02 0.015 0.16 0.25 0.0018 其余 0.217
2 0.06 0.70 2.25 0.045 0.06 0.020 0 0 0.0010 其余 0.201
3 0.08 0.40 2.06 0.085 0.04 0.050 0.20 0.10 0.0011 其余 0.218
4 0.13 0.55 1.20 0.110 0.15 0 0.16 0.25 0.0015 其余 0.183
5 0.06 0.05 1.45 0.065 0.07 0.020 0.12 0.20 0.0010 其余 0.241
6 0.10 0.15 1.90 0.095 0.09 0.008 0.15 0.22 0.0020   0.232
表2
Figure PCTCN2015070729-appb-000001
从表1和表2可以看出，本发明涉及的屈服强度890MPa级低焊接裂纹敏感性钢板的Pcm≤0.25％，屈服强度均大于890MPa，抗拉强度大于950MPa，夏氏冲击功Akv(－20℃)≥120J，板厚可达60mm，具有良好的低温韧性和焊接性。
对本发明实施例1的钢板进行焊接性能试验(小铁研试验)，在室温和50℃条件下，均未发现裂纹(见表3)，说明本发明钢种的焊接性能良好，焊接时一般不需要预热。
表3
Figure PCTCN2015070729-appb-000002
Figure PCTCN2015070729-appb-000003

Claims (3)

  1. 一种屈服强度890MPa级低焊接裂纹敏感性钢板，其化学成分重量百分比为：C 0.06～0.13wt.％、Si 0.05～0.70wt.％、Mn 1.20～2.30wt.％、Mo 0～0.25wt.％、Nb 0.03～0.11wt.％、Ti 0.002～0.050wt.％、Al0.02～0.15wt.％、B 0～0.0020wt.％，2Si+3Mn+4Mo≤8.5，其余为Fe和不可避免的杂质；且，钢板满足焊接裂纹敏感性指数Pcm≤0.25％。
  2. 一种屈服强度890MPa级低焊接裂纹敏感性钢板的制造方法，包括如下步骤：
    1)冶炼、浇铸
    按下述成分冶炼、浇铸连铸坯或钢锭，其厚度不小于成品钢板厚度的4倍；其化学成分重量百分比为：C 0.06～0.13wt.％、Si0.05～0.70wt.％、Mn 1.20～2.30wt.％、Mo 0～0.25wt.％、Nb0.03～0.11wt.％、Ti 0.002～0.050wt.％、Al 0.02～0.15wt.％、B0～0.0020wt.％，2Si+3Mn+4Mo≤8.5，其余为Fe和不可避免的杂质；
    且，钢板满足焊接裂纹敏感性指数Pcm≤0.25％；
    2)加热、轧制
    加热温度为1050～1180℃，保温时间为120～180分钟；
    轧制分为第一阶段和第二阶段轧制；
    在第一阶段轧制过程中，开轧温度为1050～1150℃，当轧件厚度到达成品钢板厚度的2～3倍时，在辊道上待温至800～860℃；
    在所述第二阶段轧制过程中，道次变形率为10～28％，终轧温度为780～840℃；
    3)冷却
    钢板以15～30℃/S的速度冷却至220～350℃，出水后空冷。
  3. 如权利要求2所述的屈服强度890MPa级低焊接裂纹敏感性钢板的制造方法，其特征是，步骤3)中，空冷采用堆垛或冷床冷却。
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