CA3180467A1 - Ultrahigh-strength dual-phase steel and manufacturing method therefor - Google Patents
Ultrahigh-strength dual-phase steel and manufacturing method therefor Download PDFInfo
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- CA3180467A1 CA3180467A1 CA3180467A CA3180467A CA3180467A1 CA 3180467 A1 CA3180467 A1 CA 3180467A1 CA 3180467 A CA3180467 A CA 3180467A CA 3180467 A CA3180467 A CA 3180467A CA 3180467 A1 CA3180467 A1 CA 3180467A1
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- 229910000885 Dual-phase steel Inorganic materials 0.000 title claims abstract description 77
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 46
- 238000005496 tempering Methods 0.000 claims abstract description 27
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 25
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 claims abstract description 23
- 229910052739 hydrogen Inorganic materials 0.000 claims abstract description 23
- 239000001257 hydrogen Substances 0.000 claims abstract description 23
- 238000005336 cracking Methods 0.000 claims abstract description 22
- 238000000137 annealing Methods 0.000 claims abstract description 21
- 230000003111 delayed effect Effects 0.000 claims abstract description 20
- 238000005097 cold rolling Methods 0.000 claims abstract description 13
- 238000009749 continuous casting Methods 0.000 claims abstract description 13
- 229910052729 chemical element Inorganic materials 0.000 claims abstract description 11
- 238000005098 hot rolling Methods 0.000 claims abstract description 9
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 9
- 238000003723 Smelting Methods 0.000 claims abstract description 5
- 239000011159 matrix material Substances 0.000 claims abstract description 4
- 238000005096 rolling process Methods 0.000 claims description 27
- VEXZGXHMUGYJMC-UHFFFAOYSA-N Hydrochloric acid Chemical compound Cl VEXZGXHMUGYJMC-UHFFFAOYSA-N 0.000 claims description 22
- 238000000034 method Methods 0.000 claims description 19
- 238000001816 cooling Methods 0.000 claims description 14
- 238000010438 heat treatment Methods 0.000 claims description 13
- 230000009467 reduction Effects 0.000 claims description 11
- 238000002791 soaking Methods 0.000 claims description 11
- 229910052717 sulfur Inorganic materials 0.000 claims description 10
- 229910052748 manganese Inorganic materials 0.000 claims description 9
- 229910052758 niobium Inorganic materials 0.000 claims description 9
- 239000012535 impurity Substances 0.000 claims description 8
- 229910052698 phosphorus Inorganic materials 0.000 claims description 8
- 230000008569 process Effects 0.000 claims description 8
- 229910052757 nitrogen Inorganic materials 0.000 claims description 7
- 150000001247 metal acetylides Chemical class 0.000 claims description 5
- 238000007747 plating Methods 0.000 claims description 4
- 238000004321 preservation Methods 0.000 claims description 4
- 229910000831 Steel Inorganic materials 0.000 description 56
- 239000010959 steel Substances 0.000 description 56
- 230000000052 comparative effect Effects 0.000 description 9
- 229910045601 alloy Inorganic materials 0.000 description 8
- 239000000956 alloy Substances 0.000 description 8
- 238000013461 design Methods 0.000 description 8
- 239000000243 solution Substances 0.000 description 7
- 230000009286 beneficial effect Effects 0.000 description 6
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 6
- 239000000203 mixture Substances 0.000 description 6
- 239000000126 substance Substances 0.000 description 6
- 229910052804 chromium Inorganic materials 0.000 description 5
- 229910052750 molybdenum Inorganic materials 0.000 description 5
- 229910052720 vanadium Inorganic materials 0.000 description 5
- 238000009792 diffusion process Methods 0.000 description 4
- 238000005246 galvanizing Methods 0.000 description 4
- 238000005728 strengthening Methods 0.000 description 4
- 229910001335 Galvanized steel Inorganic materials 0.000 description 3
- 229910001566 austenite Inorganic materials 0.000 description 3
- 230000007547 defect Effects 0.000 description 3
- 230000000694 effects Effects 0.000 description 3
- 239000008397 galvanized steel Substances 0.000 description 3
- 239000000463 material Substances 0.000 description 3
- 238000011056 performance test Methods 0.000 description 3
- 238000001556 precipitation Methods 0.000 description 3
- 239000000047 product Substances 0.000 description 3
- UQSXHKLRYXJYBZ-UHFFFAOYSA-N Iron oxide Chemical compound [Fe]=O UQSXHKLRYXJYBZ-UHFFFAOYSA-N 0.000 description 2
- 230000015572 biosynthetic process Effects 0.000 description 2
- 229910052799 carbon Inorganic materials 0.000 description 2
- 230000001427 coherent effect Effects 0.000 description 2
- 229910052742 iron Inorganic materials 0.000 description 2
- 230000007774 longterm Effects 0.000 description 2
- 239000007769 metal material Substances 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 239000002244 precipitate Substances 0.000 description 2
- 238000010583 slow cooling Methods 0.000 description 2
- 238000009864 tensile test Methods 0.000 description 2
- 230000009466 transformation Effects 0.000 description 2
- LEONUFNNVUYDNQ-UHFFFAOYSA-N vanadium atom Chemical compound [V] LEONUFNNVUYDNQ-UHFFFAOYSA-N 0.000 description 2
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 2
- INZDTEICWPZYJM-UHFFFAOYSA-N 1-(chloromethyl)-4-[4-(chloromethyl)phenyl]benzene Chemical compound C1=CC(CCl)=CC=C1C1=CC=C(CCl)C=C1 INZDTEICWPZYJM-UHFFFAOYSA-N 0.000 description 1
- 241000403513 Schizostachyum chinense Species 0.000 description 1
- QVGXLLKOCUKJST-UHFFFAOYSA-N atomic oxygen Chemical compound [O] QVGXLLKOCUKJST-UHFFFAOYSA-N 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 229910001567 cementite Inorganic materials 0.000 description 1
- 230000008859 change Effects 0.000 description 1
- 239000010960 cold rolled steel Substances 0.000 description 1
- 230000008094 contradictory effect Effects 0.000 description 1
- 230000003635 deoxygenating effect Effects 0.000 description 1
- 238000005516 engineering process Methods 0.000 description 1
- 230000006872 improvement Effects 0.000 description 1
- 239000004615 ingredient Substances 0.000 description 1
- 238000009413 insulation Methods 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 238000012986 modification Methods 0.000 description 1
- 230000004048 modification Effects 0.000 description 1
- 239000002105 nanoparticle Substances 0.000 description 1
- 230000001590 oxidative effect Effects 0.000 description 1
- 229910052760 oxygen Inorganic materials 0.000 description 1
- 239000001301 oxygen Substances 0.000 description 1
- 239000002245 particle Substances 0.000 description 1
- 238000002161 passivation Methods 0.000 description 1
- 238000012545 processing Methods 0.000 description 1
- 238000000746 purification Methods 0.000 description 1
- 230000000717 retained effect Effects 0.000 description 1
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- 239000006104 solid solution Substances 0.000 description 1
- 230000003068 static effect Effects 0.000 description 1
- 238000009628 steelmaking Methods 0.000 description 1
- 238000005482 strain hardening Methods 0.000 description 1
- 239000000758 substrate Substances 0.000 description 1
- 238000012360 testing method Methods 0.000 description 1
- 239000013585 weight reducing agent Substances 0.000 description 1
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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
- 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
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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
- C21D9/0081—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for slabs; for billets
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- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/18—Hardening; Quenching with or without subsequent tempering
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- C—CHEMISTRY; METALLURGY
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- 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
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/18—Hardening; Quenching with or without subsequent tempering
- C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
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- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/26—Methods of annealing
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- 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
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- C21D6/00—Heat treatment of ferrous alloys
- C21D6/008—Heat treatment of ferrous alloys containing Si
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- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/005—Modifying the physical properties by deformation combined with, or followed by, heat treatment of ferrous alloys
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- 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
- C21D8/021—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular fabrication or treatment of ingot or slab
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- 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
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- 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
- C21D8/0236—Cold rolling
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- 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
- C21D8/0242—Flattening; Dressing; Flexing
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- 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
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
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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
- 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
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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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
- 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
- C21D8/0278—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular surface treatment
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- 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
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C33/00—Making ferrous alloys
- C22C33/04—Making ferrous alloys by melting
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- 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
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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- 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
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- 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
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- 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
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- 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
- C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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- 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
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- C—CHEMISTRY; METALLURGY
- C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
- C25D—PROCESSES FOR THE ELECTROLYTIC OR ELECTROPHORETIC PRODUCTION OF COATINGS; ELECTROFORMING; APPARATUS THEREFOR
- C25D3/00—Electroplating: Baths therefor
- C25D3/02—Electroplating: Baths therefor from solutions
- C25D3/22—Electroplating: Baths therefor from solutions of zinc
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- C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
- C25D—PROCESSES FOR THE ELECTROLYTIC OR ELECTROPHORETIC PRODUCTION OF COATINGS; ELECTROFORMING; APPARATUS THEREFOR
- C25D5/00—Electroplating characterised by the process; Pretreatment or after-treatment of workpieces
- C25D5/34—Pretreatment of metallic surfaces to be electroplated
- C25D5/36—Pretreatment of metallic surfaces to be electroplated of iron or steel
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- 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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/004—Dispersions; Precipitations
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- 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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
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- 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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
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Abstract
Disclosed in the present invention is an ultrahigh-strength dual-phase steel. The matrix structure of the ultrahigh-strength dual-phase steel is ferrite and martensite, wherein the ferrite and the martensite are evenly distributed in an island shape. The ultrahigh-strength dual-phase steel contains the following chemical elements in percentage by mass: 0.12-0.2% of C, 0.5-1.0% of Si, 2.5-3.0% of Mn, 0.02-0.05% of Al, 0.02-0.05% of Nb, 0.02-0.05% of Ti, and 0.001-0.003% of B. Further disclosed in the present invention is a manufacturing method for the ultrahigh-strength dual-phase steel, comprising the steps of smelting and continuous casting, hot rolling, cold rolling, annealing, tempering, and leveling. The ultrahigh-strength dual-phase steel in the present invention has not only good mechanical properties but also excellent delayed cracking resistance and low initial hydrogen content, and can be suitable for manufacturing of vehicle safety structural parts.
Description
Abstract Disclosed in the present disclosure is an ultrahigh-strength dual-phase steel.
The matrix structure of the ultrahigh-strength dual-phase steel is ferrite and martensite, wherein the ferrite and the martensite are evenly distributed in an island shape. The ultrahigh-strength dual-phase steel contains the following chemical elements in percentage by mass: 0.12-0.2% of C, 0.5-1.0%
of Si, 2.5-3.0% of Mn, 0.02-0.05% of Al, 0.02-0.05% of Nb, 0.02-0.05% of Ti, and 0.001-0.003%
of B. Further disclosed in the present disclosure is a manufacturing method for the ultrahigh-strength dual-phase steel, comprising the steps of smelting and continuous casting, hot rolling, cold rolling, annealing, tempering, and leveling. The ultrahigh-strength dual-phase steel in the present disclosure has not only good mechanical properties but also excellent delayed cracking resistance and low initial hydrogen content, and can be suitable for manufacturing of vehicle safety structural parts.
ULTRAHIGH-STRENGTH DUAL-PHASE STEELAND MANUFACTURING METHOD
THE
Technical Field The present disclosure relates to a metallic material and a method of manufacturing the same, particularly to a dual-phase steel and a method of manufacturing the same.
Background Art As weight reduction and safety are required in the automotive industry, the market has an increasing demand for higher-strength steel plates. Because dual-phase steel has excellent properties such as low yield strength, high tensile strength and high initial work hardening rate in addition to its low production cost and high manufacturability, it is widely used in the production of automotive parts.
At present, strength levels of 80 kg and 100 kg are mainly demanded in the market. The current highest strength level is 1180DP grade characterized by a tensile strength of greater than or equal to 1200 MPa, a yield strength of about 850 MPa, and a total elongation of about 10%. A
critical continuous annealing process is used for the production of cold-rolled dual-phase steel whose tensile strength depends on the fraction of martensite in the annealed structure. The higher the martensite fraction, the higher the tensile strength. Thus, a higher annealing temperature is needed in the production to form a higher martensite fraction. The highest strength level of dual-phase steel that can be produced commercially is 1180 MPa, namely DP 1180 steel.
Chinese Patent Publication No. CN109504930A published on March 22, 2019 and entitled "Hot-dip Galvanized Steel Plate with Tensile Strength Greater Than 1300Mpa And Production Method Thereof" discloses a hot-dip galvanized steel plate having a tensile strength of greater than 1300 MPa and its production method. The chemical ingredients of the hot-dip galvanized steel plate substrate and their mass percentage contents are: C: 0.1-0.2%, Mn:
1.3-2.0%, S<0.005%, P<0.02%, Si: 0.2-0.3%, Als: 0.4-1.0%, Nb: 0.01-0.03%, Ti: 0.04-0.08%, B:
0.001-0.004%, Mo: 0.2-0.3%, Cr: 0.05-0.10%, V: 0.01-0.02%, and a balance of Fe and unavoidable impurities. In the slab heating step, the heating temperature is 1200-1320 C, and the heating time is 120-200 min. In the hot-rolling step, rough rolling is performed for 3-7 passes; the temperature at the finishing mill entry is 1020-1080 C; and the finishing rolling temperature is 820-880 C. The coiling temperature is 550-650 C. The production method includes steps of slab heating, hot rolling, pickling-rolling, continuous hot-dip galvanization, skin pass and passivation.
In the step of continuous hot-dip galvanization, the soaking temperature is 760-840 C; the holding time is 100-200 s; the slow cooling temperature is 680-740 C; the slow cooling rate is 10-20 C/s; the rapid cooling temperature is 420-450 C; the rapid cooling rate is 35-65 C/s; the galvanizing temperature is 458-462 C; and the galvanizing time is 5-15 s.
Chinese Patent Publication No. CN108486494A published on September 4, 2018 and entitled "Method for Producing Vanadium Microalloying 1300 Mpa Grade High-strength Hot-Rolled Steel Plate and Cold-Rolled Dual-Phase Steel Plate" discloses a method for producing vanadium microalloying 1300 MPa grade high-strength hot-rolled steel plate and cold-rolled dual-phase steel plate. The chemical composition is: 0.10-0.30 wt% C, 1.50-4.50 wt% Mn, 0.00-0.120 wt% Al, 0.00-0.90 wt% Si, 0.05-0.50% V, P<0.020 wt%, S<0.02wt%, Fe:
the balance.
The high-strength steel combines the precipitation strengthening of nano-sized vanadium carbide particles with the martensitic transformation strengthening. The strength of the existing dual-phase steel is increased significantly, and the high production efficiency is also guaranteed.
Chinese Patent Publication No. CN109628846A published on April 16, 2019 and entitled "1300 MPa Grade Ultra-high-strength Cold-rolled Steel Plate for Automobiles and Its Production Method" discloses a hot-formed steel plate and a method for manufacturing the same. The chemical composition is: C: 0.1-0.2%, Mn: 1.3-2.0%, S<0.005%, P<0.02%, Si: 0.2-0.3%, Als:
0.4-1.0%, Nb: 0.01-0.03%, Ti: 0.04-0.08%, B: 0.001-0.004%, Mo: 0.2-0.3%, Cr:
0.05-0.10%, V:
0.01-0.02%, Fe: the balance. The production method includes steps of steelmaking, continuous casting, hot rolling, pickling-rolling, continuous annealing, temper rolling and tension leveling. In the hot rolling step, the temperature for heating the slab is >1200 C; rough rolling is performed for 3-7 passes; the thickness of the intermediate slab after the rough rolling is 28-40 mm; the temperature at the finishing mill entry is 1020-1100 C; the finishing rolling temperature is 820-900 C; and the coiling temperature is 550-650 C. In the pickling-rolling step, cold rolling is performed after pickling, wherein the cold rolling reduction is >45%. In the continuous annealing step, the holding temperature in the soaking stage is 760-840 C, while the holding time is 60-225 s; and the holding temperature in the over-aging stage is 250-320 C, while the holding time in the over-aging stage is 300-1225 s.
As it can be seen, the products whose tensile strength grade is greater than or equal to 1300 MPa as disclosed by the existing patent documents are generally galvanized, and contain high Si and high Al according to some patents, which is not conducive to surface quality and production.
In some patent technologies, the products contain relatively high amounts of precious alloy elements such as Cr and Mo, and thus the production cost is high.
The matrix structure of the ultrahigh-strength dual-phase steel is ferrite and martensite, wherein the ferrite and the martensite are evenly distributed in an island shape. The ultrahigh-strength dual-phase steel contains the following chemical elements in percentage by mass: 0.12-0.2% of C, 0.5-1.0%
of Si, 2.5-3.0% of Mn, 0.02-0.05% of Al, 0.02-0.05% of Nb, 0.02-0.05% of Ti, and 0.001-0.003%
of B. Further disclosed in the present disclosure is a manufacturing method for the ultrahigh-strength dual-phase steel, comprising the steps of smelting and continuous casting, hot rolling, cold rolling, annealing, tempering, and leveling. The ultrahigh-strength dual-phase steel in the present disclosure has not only good mechanical properties but also excellent delayed cracking resistance and low initial hydrogen content, and can be suitable for manufacturing of vehicle safety structural parts.
ULTRAHIGH-STRENGTH DUAL-PHASE STEELAND MANUFACTURING METHOD
THE
Technical Field The present disclosure relates to a metallic material and a method of manufacturing the same, particularly to a dual-phase steel and a method of manufacturing the same.
Background Art As weight reduction and safety are required in the automotive industry, the market has an increasing demand for higher-strength steel plates. Because dual-phase steel has excellent properties such as low yield strength, high tensile strength and high initial work hardening rate in addition to its low production cost and high manufacturability, it is widely used in the production of automotive parts.
At present, strength levels of 80 kg and 100 kg are mainly demanded in the market. The current highest strength level is 1180DP grade characterized by a tensile strength of greater than or equal to 1200 MPa, a yield strength of about 850 MPa, and a total elongation of about 10%. A
critical continuous annealing process is used for the production of cold-rolled dual-phase steel whose tensile strength depends on the fraction of martensite in the annealed structure. The higher the martensite fraction, the higher the tensile strength. Thus, a higher annealing temperature is needed in the production to form a higher martensite fraction. The highest strength level of dual-phase steel that can be produced commercially is 1180 MPa, namely DP 1180 steel.
Chinese Patent Publication No. CN109504930A published on March 22, 2019 and entitled "Hot-dip Galvanized Steel Plate with Tensile Strength Greater Than 1300Mpa And Production Method Thereof" discloses a hot-dip galvanized steel plate having a tensile strength of greater than 1300 MPa and its production method. The chemical ingredients of the hot-dip galvanized steel plate substrate and their mass percentage contents are: C: 0.1-0.2%, Mn:
1.3-2.0%, S<0.005%, P<0.02%, Si: 0.2-0.3%, Als: 0.4-1.0%, Nb: 0.01-0.03%, Ti: 0.04-0.08%, B:
0.001-0.004%, Mo: 0.2-0.3%, Cr: 0.05-0.10%, V: 0.01-0.02%, and a balance of Fe and unavoidable impurities. In the slab heating step, the heating temperature is 1200-1320 C, and the heating time is 120-200 min. In the hot-rolling step, rough rolling is performed for 3-7 passes; the temperature at the finishing mill entry is 1020-1080 C; and the finishing rolling temperature is 820-880 C. The coiling temperature is 550-650 C. The production method includes steps of slab heating, hot rolling, pickling-rolling, continuous hot-dip galvanization, skin pass and passivation.
In the step of continuous hot-dip galvanization, the soaking temperature is 760-840 C; the holding time is 100-200 s; the slow cooling temperature is 680-740 C; the slow cooling rate is 10-20 C/s; the rapid cooling temperature is 420-450 C; the rapid cooling rate is 35-65 C/s; the galvanizing temperature is 458-462 C; and the galvanizing time is 5-15 s.
Chinese Patent Publication No. CN108486494A published on September 4, 2018 and entitled "Method for Producing Vanadium Microalloying 1300 Mpa Grade High-strength Hot-Rolled Steel Plate and Cold-Rolled Dual-Phase Steel Plate" discloses a method for producing vanadium microalloying 1300 MPa grade high-strength hot-rolled steel plate and cold-rolled dual-phase steel plate. The chemical composition is: 0.10-0.30 wt% C, 1.50-4.50 wt% Mn, 0.00-0.120 wt% Al, 0.00-0.90 wt% Si, 0.05-0.50% V, P<0.020 wt%, S<0.02wt%, Fe:
the balance.
The high-strength steel combines the precipitation strengthening of nano-sized vanadium carbide particles with the martensitic transformation strengthening. The strength of the existing dual-phase steel is increased significantly, and the high production efficiency is also guaranteed.
Chinese Patent Publication No. CN109628846A published on April 16, 2019 and entitled "1300 MPa Grade Ultra-high-strength Cold-rolled Steel Plate for Automobiles and Its Production Method" discloses a hot-formed steel plate and a method for manufacturing the same. The chemical composition is: C: 0.1-0.2%, Mn: 1.3-2.0%, S<0.005%, P<0.02%, Si: 0.2-0.3%, Als:
0.4-1.0%, Nb: 0.01-0.03%, Ti: 0.04-0.08%, B: 0.001-0.004%, Mo: 0.2-0.3%, Cr:
0.05-0.10%, V:
0.01-0.02%, Fe: the balance. The production method includes steps of steelmaking, continuous casting, hot rolling, pickling-rolling, continuous annealing, temper rolling and tension leveling. In the hot rolling step, the temperature for heating the slab is >1200 C; rough rolling is performed for 3-7 passes; the thickness of the intermediate slab after the rough rolling is 28-40 mm; the temperature at the finishing mill entry is 1020-1100 C; the finishing rolling temperature is 820-900 C; and the coiling temperature is 550-650 C. In the pickling-rolling step, cold rolling is performed after pickling, wherein the cold rolling reduction is >45%. In the continuous annealing step, the holding temperature in the soaking stage is 760-840 C, while the holding time is 60-225 s; and the holding temperature in the over-aging stage is 250-320 C, while the holding time in the over-aging stage is 300-1225 s.
As it can be seen, the products whose tensile strength grade is greater than or equal to 1300 MPa as disclosed by the existing patent documents are generally galvanized, and contain high Si and high Al according to some patents, which is not conducive to surface quality and production.
In some patent technologies, the products contain relatively high amounts of precious alloy elements such as Cr and Mo, and thus the production cost is high.
2 Summary One of the objects of the present disclosure is to provide an ultra-high-strength dual-phase steel. By a reasonable design of the chemical elements in the ultra-high-strength dual-phase steel, i.e. a design of medium Si and low Al to reduce the use of alloy elements such as Si and Al, the problems with the surface quality caused by high Si and the slab defects caused by high Al are avoided.
In addition, precious alloy elements such as Cr and Mo are not used in the ultra-high-strength dual-phase steel according to the present disclosure, and thus the alloy cost is controlled effectively. At the same time, the contents of impurity elements P
and S are reduced, which is beneficial to promotion of performances and improvement of delayed cracking. The ultra-high-strength dual-phase steel has a yield strength of >900 MPa, preferably>930 MPa, a tensile strength of >1300 MPa, preferably >1320 MPa, an elongation after fracture of >5%, preferably >5.5%, an initial hydrogen content of <10 ppm, preferably < 7 ppm;
and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength, and preferably does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength. It can be used effectively for manufacture of automotive safety structural parts. It is highly valuable and promising for popularization and application.
In order to achieve the above object, the present disclosure provides an ultra-high-strength dual-phase steel having a matrix structure of ferrite + martensite, wherein ferrite and martensite are distributed evenly like islands, and wherein the ultra-high-strength dual-phase steel comprises the following chemical elements in mass percentages, in addition to Fe:
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%-0.003%.
Further, the ultra-high-strength dual-phase steel comprises the following chemical elements in mass percentages:
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%-0.003%, and a balance of Fe and other unavoidable impurities.
In the ultra-high-strength dual-phase steel, the various chemical elements are designed according to the following principles:
C: In the ultra-high-strength dual-phase steel according to the present disclosure, C is a solid solution strengthening element, and it is a guarantee for the material to obtain high strength.
However, it should be noted that the higher the C content in the steel, the harder the martensite
In addition, precious alloy elements such as Cr and Mo are not used in the ultra-high-strength dual-phase steel according to the present disclosure, and thus the alloy cost is controlled effectively. At the same time, the contents of impurity elements P
and S are reduced, which is beneficial to promotion of performances and improvement of delayed cracking. The ultra-high-strength dual-phase steel has a yield strength of >900 MPa, preferably>930 MPa, a tensile strength of >1300 MPa, preferably >1320 MPa, an elongation after fracture of >5%, preferably >5.5%, an initial hydrogen content of <10 ppm, preferably < 7 ppm;
and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength, and preferably does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength. It can be used effectively for manufacture of automotive safety structural parts. It is highly valuable and promising for popularization and application.
In order to achieve the above object, the present disclosure provides an ultra-high-strength dual-phase steel having a matrix structure of ferrite + martensite, wherein ferrite and martensite are distributed evenly like islands, and wherein the ultra-high-strength dual-phase steel comprises the following chemical elements in mass percentages, in addition to Fe:
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%-0.003%.
Further, the ultra-high-strength dual-phase steel comprises the following chemical elements in mass percentages:
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%-0.003%, and a balance of Fe and other unavoidable impurities.
In the ultra-high-strength dual-phase steel, the various chemical elements are designed according to the following principles:
C: In the ultra-high-strength dual-phase steel according to the present disclosure, C is a solid solution strengthening element, and it is a guarantee for the material to obtain high strength.
However, it should be noted that the higher the C content in the steel, the harder the martensite
3 and the greater the tendency for delayed cracking to occur. Therefore, when a product is designed, it's better to choose a low-carbon design. In the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of C is controlled at 0.12-0.2%.
In some preferred embodiments, the mass percentage of C may be controlled at 0.14-0.18%.
Si: In the ultra-high-strength dual-phase steel according to the present disclosure, Si has an effect of increasing the elongation of the steel. Si also has a great influence on the structure of the steel. Particularly, it promotes purification of ferrite and formation of retained austenite. At the same time, it can improve the tempering resistance of martensite, and inhibit precipitation and growth of Fe3C, so that the dominated precipitates formed during tempering are E carbides.
However, it should be noted that when the mass percentage of Si in the steel is less than 0.5%, the elongation and tempering resistance of the steel will be affected; if the mass percentage of Si is higher than 1.0%, other metallurgical quality defects will be caused.
Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of Si is controlled at 0.5-1.0%.
Mn: In the ultra-high-strength dual-phase steel according to the present disclosure, Mn is an element that strongly improves the hardenability of austenite, and it can improve the strength of the steel effectively by forming more martensite. Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of Mn is controlled at 2.5-3.0%.
In some preferred embodiments, the mass percentage of Mn may be controlled at 2.5-2.8%.
Al: In the ultra-high strength dual-phase steel according to the present disclosure, Al is a deoxygenating element. It can remove oxygen and refine grains in the steel.
Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of Al is controlled at 0.02-0.05%.
Nb and Ti: In the ultra-high-strength dual-phase steel according to the present disclosure, Nb and Ti are elements for precipitation of carbonitrides. They can refine grains, precipitate carbonitrides, and improve the strength of the material. They can be added separately or in combination. However, it should be noted that if the mass percentage of Nb or Ti in the steel is higher than 0.05%, the strengthening effect is not obvious. Therefore, in the ultra-high-strength dual-phase steel of the present disclosure, the mass percentage of Nb is controlled at 0.02-0.05%, and the mass percentage of Ti is controlled at 0.02-0.05%.
B: In the ultra-high-strength dual-phase steel according to the present disclosure, B is used as a strong element for hardenability. An appropriate amount of B can increase the hardenability of the steel, and promote formation of martensite. Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of B is controlled at
In some preferred embodiments, the mass percentage of C may be controlled at 0.14-0.18%.
Si: In the ultra-high-strength dual-phase steel according to the present disclosure, Si has an effect of increasing the elongation of the steel. Si also has a great influence on the structure of the steel. Particularly, it promotes purification of ferrite and formation of retained austenite. At the same time, it can improve the tempering resistance of martensite, and inhibit precipitation and growth of Fe3C, so that the dominated precipitates formed during tempering are E carbides.
However, it should be noted that when the mass percentage of Si in the steel is less than 0.5%, the elongation and tempering resistance of the steel will be affected; if the mass percentage of Si is higher than 1.0%, other metallurgical quality defects will be caused.
Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of Si is controlled at 0.5-1.0%.
Mn: In the ultra-high-strength dual-phase steel according to the present disclosure, Mn is an element that strongly improves the hardenability of austenite, and it can improve the strength of the steel effectively by forming more martensite. Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of Mn is controlled at 2.5-3.0%.
In some preferred embodiments, the mass percentage of Mn may be controlled at 2.5-2.8%.
Al: In the ultra-high strength dual-phase steel according to the present disclosure, Al is a deoxygenating element. It can remove oxygen and refine grains in the steel.
Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of Al is controlled at 0.02-0.05%.
Nb and Ti: In the ultra-high-strength dual-phase steel according to the present disclosure, Nb and Ti are elements for precipitation of carbonitrides. They can refine grains, precipitate carbonitrides, and improve the strength of the material. They can be added separately or in combination. However, it should be noted that if the mass percentage of Nb or Ti in the steel is higher than 0.05%, the strengthening effect is not obvious. Therefore, in the ultra-high-strength dual-phase steel of the present disclosure, the mass percentage of Nb is controlled at 0.02-0.05%, and the mass percentage of Ti is controlled at 0.02-0.05%.
B: In the ultra-high-strength dual-phase steel according to the present disclosure, B is used as a strong element for hardenability. An appropriate amount of B can increase the hardenability of the steel, and promote formation of martensite. Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of B is controlled at
4 0.001%-0.003%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, the unavoidable impurities include the P, S and N elements, and the contents thereof are controlled to be at least one of the following: P <0.01%, S<0.002%, N<0.004%.
In the above technical solution, in the ultra-high-strength dual-phase steel according to the present disclosure, the P, S and N elements are all unavoidable impurity elements in the steel. It's better to lower the contents of the P, S and N elements in the steel as far as possible. S tends to form MnS inclusions which will seriously affect the hole expansion rate. The P
element may reduce the toughness of the steel, which is not conducive to the delayed cracking performance. An unduly high content of the N element in the steel is prone to causing cracks on the surface of the slab, which will greatly affect the performances of the steel. Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of P
is controlled at P<0.01%; the mass percentage of S is controlled at S<0.002%; and the mass percentage of N is controlled at N <0.004%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage contents of the chemical elements satisfy at least one of the following:
C: 0.14-0.18%, Mn: 2.5-2.8%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, the phase proportion (by volume) of martensite is >90%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, martensite further comprises coherently distributed E carbides.
Further, the performances of the ultra-high-strength dual-phase steel according to the present disclosure meet at least one of the following: yield strength >900 MPa, tensile strength >1300 MPa, elongation after fracture >5%, initial hydrogen content <10 ppm; no delayed cracking when soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength.
Further, the performances of the ultra-high-strength dual-phase steel according to the present disclosure meet at least one of the following: yield strength >930 MPa, tensile strength >1320 MPa, elongation after fracture >5.5%, initial hydrogen content <7 ppm; no delayed cracking when soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength.
Further, the ultra-high-strength dual-phase steel according to the present disclosure has a yield strength of >930 MPa, a tensile strength of >1320 MPa, an elongation after fracture of 25.5%, an initial hydrogen content of <7 ppm; and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength.
Accordingly, another object of the present disclosure is to provide a method for manufacturing an ultra-high-strength dual-phase steel. The ultra-high-strength dual-phase steel manufactured by this method has a yield strength of 2900 MPa, a tensile strength of 21300 MPa, an elongation after fracture of 25%, an initial hydrogen content of <10 ppm, and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength. It can be used effectively for manufacture of automotive safety structural parts. It is highly valuable and promising for popularization and application.
To fulfil the above object, the present disclosure proposes a method for manufacturing the above ultra-high-strength dual-phase steel, comprising steps of:
(1) Smelting and continuous casting;
(2) Hot rolling;
(3) Cold rolling;
(4) Annealing: heating to an annealing soaking temperature of 800-850 C, preferably 805-845 C at a heating rate of 3-10 C/s, the annealing time being 40-200 s;
and then rapidly cooling at a rate of 30-80 C/s, a starting temperature of the rapid cooling being 670-730 C;
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, the unavoidable impurities include the P, S and N elements, and the contents thereof are controlled to be at least one of the following: P <0.01%, S<0.002%, N<0.004%.
In the above technical solution, in the ultra-high-strength dual-phase steel according to the present disclosure, the P, S and N elements are all unavoidable impurity elements in the steel. It's better to lower the contents of the P, S and N elements in the steel as far as possible. S tends to form MnS inclusions which will seriously affect the hole expansion rate. The P
element may reduce the toughness of the steel, which is not conducive to the delayed cracking performance. An unduly high content of the N element in the steel is prone to causing cracks on the surface of the slab, which will greatly affect the performances of the steel. Therefore, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage of P
is controlled at P<0.01%; the mass percentage of S is controlled at S<0.002%; and the mass percentage of N is controlled at N <0.004%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, the mass percentage contents of the chemical elements satisfy at least one of the following:
C: 0.14-0.18%, Mn: 2.5-2.8%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, the phase proportion (by volume) of martensite is >90%.
Further, in the ultra-high-strength dual-phase steel according to the present disclosure, martensite further comprises coherently distributed E carbides.
Further, the performances of the ultra-high-strength dual-phase steel according to the present disclosure meet at least one of the following: yield strength >900 MPa, tensile strength >1300 MPa, elongation after fracture >5%, initial hydrogen content <10 ppm; no delayed cracking when soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength.
Further, the performances of the ultra-high-strength dual-phase steel according to the present disclosure meet at least one of the following: yield strength >930 MPa, tensile strength >1320 MPa, elongation after fracture >5.5%, initial hydrogen content <7 ppm; no delayed cracking when soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength.
Further, the ultra-high-strength dual-phase steel according to the present disclosure has a yield strength of >930 MPa, a tensile strength of >1320 MPa, an elongation after fracture of 25.5%, an initial hydrogen content of <7 ppm; and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength.
Accordingly, another object of the present disclosure is to provide a method for manufacturing an ultra-high-strength dual-phase steel. The ultra-high-strength dual-phase steel manufactured by this method has a yield strength of 2900 MPa, a tensile strength of 21300 MPa, an elongation after fracture of 25%, an initial hydrogen content of <10 ppm, and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength. It can be used effectively for manufacture of automotive safety structural parts. It is highly valuable and promising for popularization and application.
To fulfil the above object, the present disclosure proposes a method for manufacturing the above ultra-high-strength dual-phase steel, comprising steps of:
(1) Smelting and continuous casting;
(2) Hot rolling;
(3) Cold rolling;
(4) Annealing: heating to an annealing soaking temperature of 800-850 C, preferably 805-845 C at a heating rate of 3-10 C/s, the annealing time being 40-200 s;
and then rapidly cooling at a rate of 30-80 C/s, a starting temperature of the rapid cooling being 670-730 C;
(5) Tempering: tempering temperature: 260-320 C, preferably 260-310 C;
tempering time:
100-400 s, preferably 100-300 s;
tempering time:
100-400 s, preferably 100-300 s;
(6) Temper rolling;
(7) Electro-galvanizing.
A combination of high temperature soaking and medium temperature tempering is employed for the annealing. The high temperature soaking gives rise to more austenite transformation, and thus more martensite is obtained during the subsequent rapid cooling, which finally guarantees higher strength before tempering. The medium temperature tempering provides the material with a moderate yield ratio on the one hand, and on the other hand, it has a better effect in improving the delayed cracking performance. In a preferred embodiment, the ultra-high-strength dual-phase steel according to the present disclosure has a yield ratio of 0.70-0.75.
In the method for manufacturing the ultra-high strength dual-phase steel according to the present disclosure, by adopting the medium to low temperature tempering treatment after the continuous annealing and controlling the relevant process parameters, uniform, fine and dispersive coherent E carbides can be easily precipitated during tempering of martensite on the one hand, and on the other hand, the manner of long-term tempering at medium to low temperature enables removal of superfluous hydrogen from the steel plate to the greatest extent, i.e. diffusion of it out of the steel plate, so that the amount of hydrogen in its original state in the steel plate can be reduced. This is not only beneficial to reduce the hardness of martensite and the diffusion of hydrogen inside the steel plate, but also very beneficial to the mechanical properties and delayed cracking performance of the steel.
Further, in the manufacturing method according to the present disclosure, in step (1), a drawing speed in the continuous casting is controlled at 0.9-1.5 m/min.
In the above technical solution, in the manufacturing method according to the present disclosure, in step (1), the continuous casting may be performed in a secondary cooling mode with a large amount of water to minimize segregation.
Further, in the manufacturing method according to the present disclosure, in step (2), the cast slab is controlled to be soaked at a temperature of 1220-1260 C, preferably 1220-1250 C; then rolled with a finishing rolling temperature being controlled at 880-920 C;
then cooled at a rate of 20-70 C/s after rolling; then coiled at a coiling temperature of 600-650 C, preferably 605-645 C; and then subjected to heat preservation treatment after coiling.
Preferably, the heat preservation treatment is performed for 1-5 hours after coiling.
In the method for manufacturing the ultra-high-strength dual-phase steel according to the present disclosure, in step (2), in order to guarantee the stability of the rolling load, the heating temperature is controlled at 1220 C or higher. Meanwhile, the upper limit of the heating temperature is controlled to be 1260 C in order to prevent increase of oxidative burning loss.
Therefore, the cast slab is finally controlled to be soaked at a temperature of 1220-1260 C.
Further, in the manufacturing method according to the present disclosure, in step (3), the cold rolling reduction rate is controlled at 45-65%.
In the above technical solution, in the step (3), before cold rolling at a cold rolling reduction rate controlled at 45-65%, iron oxide scale on the surface of the steel plate can be removed by pickling.
Further, in the manufacturing method according to the present disclosure, in step (6), the temper rolling reduction rate is controlled at <0.3%.
In the above technical solution according to the present disclosure, in step (6), in order to guarantee the flatness of the steel plate, a certain amount of temper rolling needs to be performed, but an excessively large amount of temper rolling will increase the yield strength of the steel too much. Therefore, in the manufacturing method according to the present disclosure, the temper rolling reduction rate is controlled at <0.3%.
In the above technical solution according to the present disclosure, step (7) may be performed by a conventional electro-galvanizing method. Preferably, double-side plating is performed, and the weight of the plating layer on one side is in the range of 10-100 g/m2.
Compared with the prior art, the ultra-high-strength dual-phase steel and the manufacturing method thereof have the following advantages and beneficial effects:
The composition of the ultra-high-strength dual-phase steel according to the present disclosure is designed reasonably. That is, a design of medium Si and low Al is employed to reduce the use of alloy elements such as Si and Al, so that the problems with the surface quality caused by high Si and the slab defects caused by high Al are avoided. In addition, the steel is free of precious alloy elements such as Cr and Mo, and the alloy content is low.
The manufacturability and economic efficiency of the steel are very good. The alloy cost is controlled effectively. The ultra-high-strength dual-phase steel has a yield strength of 2900 MPa, a tensile strength of 21300 MPa, an elongation after fracture of >5%, an initial hydrogen content of <10 ppm; and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength. It can be used effectively for manufacture of automotive safety structural parts. It is highly valuable and promising for popularization and application.
In addition, in the manufacturing method according to the present disclosure, by adopting the medium to low temperature tempering treatment after the continuous annealing and controlling the relevant process parameters, uniform, fine and dispersive coherent E
carbides can be easily precipitated during tempering of martensite on the one hand, and on the other hand, the manner of long-term tempering at medium to low temperature enables removal of superfluous hydrogen from the steel plate to the greatest extent, i.e. diffusion of it out of the steel plate, so that the amount of hydrogen in its original state in the steel plate can be reduced.
This is not only beneficial to reduce the hardness of martensite and the diffusion of hydrogen inside the steel plate, but also very beneficial to the mechanical properties and delayed cracking performance of the steel, thereby effectively ensuring that the produced ultra-high strength dual-phase steel has excellent mechanical properties, excellent resistance to delayed cracking, and a lower initial hydrogen content.
Description of the Drawing Figure 1 shows the structure of the cold-rolled and annealed dual-phase steel of Example 1.
Detailed Description
A combination of high temperature soaking and medium temperature tempering is employed for the annealing. The high temperature soaking gives rise to more austenite transformation, and thus more martensite is obtained during the subsequent rapid cooling, which finally guarantees higher strength before tempering. The medium temperature tempering provides the material with a moderate yield ratio on the one hand, and on the other hand, it has a better effect in improving the delayed cracking performance. In a preferred embodiment, the ultra-high-strength dual-phase steel according to the present disclosure has a yield ratio of 0.70-0.75.
In the method for manufacturing the ultra-high strength dual-phase steel according to the present disclosure, by adopting the medium to low temperature tempering treatment after the continuous annealing and controlling the relevant process parameters, uniform, fine and dispersive coherent E carbides can be easily precipitated during tempering of martensite on the one hand, and on the other hand, the manner of long-term tempering at medium to low temperature enables removal of superfluous hydrogen from the steel plate to the greatest extent, i.e. diffusion of it out of the steel plate, so that the amount of hydrogen in its original state in the steel plate can be reduced. This is not only beneficial to reduce the hardness of martensite and the diffusion of hydrogen inside the steel plate, but also very beneficial to the mechanical properties and delayed cracking performance of the steel.
Further, in the manufacturing method according to the present disclosure, in step (1), a drawing speed in the continuous casting is controlled at 0.9-1.5 m/min.
In the above technical solution, in the manufacturing method according to the present disclosure, in step (1), the continuous casting may be performed in a secondary cooling mode with a large amount of water to minimize segregation.
Further, in the manufacturing method according to the present disclosure, in step (2), the cast slab is controlled to be soaked at a temperature of 1220-1260 C, preferably 1220-1250 C; then rolled with a finishing rolling temperature being controlled at 880-920 C;
then cooled at a rate of 20-70 C/s after rolling; then coiled at a coiling temperature of 600-650 C, preferably 605-645 C; and then subjected to heat preservation treatment after coiling.
Preferably, the heat preservation treatment is performed for 1-5 hours after coiling.
In the method for manufacturing the ultra-high-strength dual-phase steel according to the present disclosure, in step (2), in order to guarantee the stability of the rolling load, the heating temperature is controlled at 1220 C or higher. Meanwhile, the upper limit of the heating temperature is controlled to be 1260 C in order to prevent increase of oxidative burning loss.
Therefore, the cast slab is finally controlled to be soaked at a temperature of 1220-1260 C.
Further, in the manufacturing method according to the present disclosure, in step (3), the cold rolling reduction rate is controlled at 45-65%.
In the above technical solution, in the step (3), before cold rolling at a cold rolling reduction rate controlled at 45-65%, iron oxide scale on the surface of the steel plate can be removed by pickling.
Further, in the manufacturing method according to the present disclosure, in step (6), the temper rolling reduction rate is controlled at <0.3%.
In the above technical solution according to the present disclosure, in step (6), in order to guarantee the flatness of the steel plate, a certain amount of temper rolling needs to be performed, but an excessively large amount of temper rolling will increase the yield strength of the steel too much. Therefore, in the manufacturing method according to the present disclosure, the temper rolling reduction rate is controlled at <0.3%.
In the above technical solution according to the present disclosure, step (7) may be performed by a conventional electro-galvanizing method. Preferably, double-side plating is performed, and the weight of the plating layer on one side is in the range of 10-100 g/m2.
Compared with the prior art, the ultra-high-strength dual-phase steel and the manufacturing method thereof have the following advantages and beneficial effects:
The composition of the ultra-high-strength dual-phase steel according to the present disclosure is designed reasonably. That is, a design of medium Si and low Al is employed to reduce the use of alloy elements such as Si and Al, so that the problems with the surface quality caused by high Si and the slab defects caused by high Al are avoided. In addition, the steel is free of precious alloy elements such as Cr and Mo, and the alloy content is low.
The manufacturability and economic efficiency of the steel are very good. The alloy cost is controlled effectively. The ultra-high-strength dual-phase steel has a yield strength of 2900 MPa, a tensile strength of 21300 MPa, an elongation after fracture of >5%, an initial hydrogen content of <10 ppm; and it does not experience delayed cracking when it is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength. It can be used effectively for manufacture of automotive safety structural parts. It is highly valuable and promising for popularization and application.
In addition, in the manufacturing method according to the present disclosure, by adopting the medium to low temperature tempering treatment after the continuous annealing and controlling the relevant process parameters, uniform, fine and dispersive coherent E
carbides can be easily precipitated during tempering of martensite on the one hand, and on the other hand, the manner of long-term tempering at medium to low temperature enables removal of superfluous hydrogen from the steel plate to the greatest extent, i.e. diffusion of it out of the steel plate, so that the amount of hydrogen in its original state in the steel plate can be reduced.
This is not only beneficial to reduce the hardness of martensite and the diffusion of hydrogen inside the steel plate, but also very beneficial to the mechanical properties and delayed cracking performance of the steel, thereby effectively ensuring that the produced ultra-high strength dual-phase steel has excellent mechanical properties, excellent resistance to delayed cracking, and a lower initial hydrogen content.
Description of the Drawing Figure 1 shows the structure of the cold-rolled and annealed dual-phase steel of Example 1.
Detailed Description
8 The ultra-high-strength dual-phase steel and the method for manufacturing the same according to the disclosure will be further explained and illustrated with reference to the specific Examples. Nonetheless, the explanation and illustration are not intended to unduly limit the technical solution of the disclosure.
Examples 1-7 and Comparative Examples 1-14 Table 1 lists the mass percentages of various chemical elements in the steel grades corresponding to the ultra-high-strength dual-phase steels in Examples 1-7 and the steels in Comparative Examples 1-14.
Table 1 (wt%, the balance is Fe and other unavoidable impurities except for P, S and N) Steel grade C Si Mn P S Nb Ti Al N B
Ex. 1 A 0.122 0.54 2.52 0.01 0.001 0.036 0.033 0.034 0.0030 0.0011 Ex. 2 B 0.147 0.72 2.64 0.008 0.0008 0.044 0.038 0.043 0.0027 0.0024 Ex. 3 C 0.131 0.63 2.98 0.009 0.002 0.022 0.045 0.028 0.0033 0.0029 Ex. 4 D 0.158 0.84 2.56 0.007 0.0007 0.043 0.037 0.022 0.0028 0.0015 Ex. 5 E 0.164 0.55 2.82 0.01 0.001 0.027 0.025 0.038 0.0032 0.0018 Ex. 6 F 0.173 0.95 2.73 0.005 0.0015 0.033 0.038 0.034 0.0029 0.0016 Ex. 7 G 0.194 0.66 2.65 0.009 0.0009 0.048 0.04 0.049 0.0026 0.0021 Comp. Ex. 1 H 0.118 0.67 2.65 0.006 0.002 0.025 0.043 0.034 0.0028 0.0013 Comp. Ex. 2 I 0.205 0.58 2.53 0.009 0.0008 0.038 0.025 0.027 0.0034 0.0025 Comp. Ex. 3 j 0.144 0.47 2.42 0.01 0.0016 0.043 0.035 0.022 0.0037 0.0016 Comp. Ex. 4 K
0.153 0.65 3.05 0.0007 0.0013 0.037 0.043 0.037 0.0028 0.0022 Comp. Ex. 5 L 0.149 0.72 2.75 0.01 0.0008 0.025 0.013 0.025 0.0022 0.0017 Comp. Ex. 6 ivi 0.162 0.64 2.58 0.0005 0.0012 0.01 0.035 0.039 0.0032 0.0024 Comp. Ex. 7-14 N 0.175 0.77 2.55 0.008 0.0009 0.033 0.026 0.046 0.0028 0.0018 The ultra-high-strength dual-phase steels in Examples 1-7 according to the present disclosure and the steels in Comparative Examples 1-14 were all prepared by the following steps:
(1) Smelting and continuous casting: The drawing speed in the continuous casting was controlled to be 0.9-1.5 m/min during the continuous casting process, and the continuous casting was carried out in a secondary cooling mode with a large amount of water;
(2) Hot rolling: The cast slab was soaked at a temperature controlled at 1220-1260 C, and then rolled, wherein the finishing rolling temperature was controlled at 880-920 C. After rolling, the steel was cooled at a rate of 20-70 C/s. Then, the steel was coiled at a coiling temperature of 600-650 C. After coiling, an insulation cover was used to perform heat preservation treatment;
Examples 1-7 and Comparative Examples 1-14 Table 1 lists the mass percentages of various chemical elements in the steel grades corresponding to the ultra-high-strength dual-phase steels in Examples 1-7 and the steels in Comparative Examples 1-14.
Table 1 (wt%, the balance is Fe and other unavoidable impurities except for P, S and N) Steel grade C Si Mn P S Nb Ti Al N B
Ex. 1 A 0.122 0.54 2.52 0.01 0.001 0.036 0.033 0.034 0.0030 0.0011 Ex. 2 B 0.147 0.72 2.64 0.008 0.0008 0.044 0.038 0.043 0.0027 0.0024 Ex. 3 C 0.131 0.63 2.98 0.009 0.002 0.022 0.045 0.028 0.0033 0.0029 Ex. 4 D 0.158 0.84 2.56 0.007 0.0007 0.043 0.037 0.022 0.0028 0.0015 Ex. 5 E 0.164 0.55 2.82 0.01 0.001 0.027 0.025 0.038 0.0032 0.0018 Ex. 6 F 0.173 0.95 2.73 0.005 0.0015 0.033 0.038 0.034 0.0029 0.0016 Ex. 7 G 0.194 0.66 2.65 0.009 0.0009 0.048 0.04 0.049 0.0026 0.0021 Comp. Ex. 1 H 0.118 0.67 2.65 0.006 0.002 0.025 0.043 0.034 0.0028 0.0013 Comp. Ex. 2 I 0.205 0.58 2.53 0.009 0.0008 0.038 0.025 0.027 0.0034 0.0025 Comp. Ex. 3 j 0.144 0.47 2.42 0.01 0.0016 0.043 0.035 0.022 0.0037 0.0016 Comp. Ex. 4 K
0.153 0.65 3.05 0.0007 0.0013 0.037 0.043 0.037 0.0028 0.0022 Comp. Ex. 5 L 0.149 0.72 2.75 0.01 0.0008 0.025 0.013 0.025 0.0022 0.0017 Comp. Ex. 6 ivi 0.162 0.64 2.58 0.0005 0.0012 0.01 0.035 0.039 0.0032 0.0024 Comp. Ex. 7-14 N 0.175 0.77 2.55 0.008 0.0009 0.033 0.026 0.046 0.0028 0.0018 The ultra-high-strength dual-phase steels in Examples 1-7 according to the present disclosure and the steels in Comparative Examples 1-14 were all prepared by the following steps:
(1) Smelting and continuous casting: The drawing speed in the continuous casting was controlled to be 0.9-1.5 m/min during the continuous casting process, and the continuous casting was carried out in a secondary cooling mode with a large amount of water;
(2) Hot rolling: The cast slab was soaked at a temperature controlled at 1220-1260 C, and then rolled, wherein the finishing rolling temperature was controlled at 880-920 C. After rolling, the steel was cooled at a rate of 20-70 C/s. Then, the steel was coiled at a coiling temperature of 600-650 C. After coiling, an insulation cover was used to perform heat preservation treatment;
9 (3) Cold rolling: The cold rolling reduction rate was controlled at 45-65%;
(4) Annealing: The temperature was raised to the annealing soaking temperature of 800-850 C at a heating rate of 3-10 C/s, wherein the annealing time was 40-200 s.
Then, rapid cooling was performed at a rate of 30-80 C/s, wherein the starting temperature of the rapid cooling was 670-730 C;
(5) Tempering: The tempering temperature was 260-320 C, and the tempering time was 100-400 s;
(6) Temper rolling: The temper rolling reduction rate was controlled at <0.3%;
(7) Double-side electro-galvanization: The weight of the plating layer on each side was
(4) Annealing: The temperature was raised to the annealing soaking temperature of 800-850 C at a heating rate of 3-10 C/s, wherein the annealing time was 40-200 s.
Then, rapid cooling was performed at a rate of 30-80 C/s, wherein the starting temperature of the rapid cooling was 670-730 C;
(5) Tempering: The tempering temperature was 260-320 C, and the tempering time was 100-400 s;
(6) Temper rolling: The temper rolling reduction rate was controlled at <0.3%;
(7) Double-side electro-galvanization: The weight of the plating layer on each side was
10-100 g/m2.
It should be noted that the chemical compositions of the ultra-high-strength dual-phase steel in Examples 1-7 and the related process parameters all met the control requirements of the design specification according to the present disclosure. The chemical compositions of the steels in Comparative Examples 1-6 all included parameters that failed to meet the requirements of the design according to the present disclosure. Although the chemical composition of steel grade N in Comparative Examples 7-14 met the requirements of the design according to the present disclosure, the related process parameters all included parameters that failed to meet the requirements of the design according to the present disclosure.
Tables 2-1 and 2-2 list the specific process parameters for the ultra-high-strength dual-phase steels in Examples 1-7 and the steels in Comparative Examples 1-14.
Table 2-1 Step (1) Step (2) Step (3) Drawing speed Finishing Steel Soaking Cooling Coiling Cold rolling No. in continuous rolling grade temperature rate temperature reduction casting temperature ( C) ( C/s) ( C) rate (%) (m/min) ( C) Ex. 1 A 1.0 1250 895 25 605 Ex. 2 B 1.2 1245 880 30 625 Ex. 3 c 1.5 1220 890 45 645 Ex. 4 D 0.9 1234 905 50 625 Ex. 5 E 1.1 1224 910 35 615 Ex. 6 F 1.3 1240 890 60 600 Ex. 7 G 1.0 1250 885 65 630 Comp. H 1.2 1230 895 40 615 Ex. 1 Comp.
I 0.9 1228 915 55 620 Ex. 2 Comp.
J 1.4 1250 890 70 625 Ex. 3 Comp.
K 1.3 1255 900 45 615 Ex. 4 Comp.
L 0.6 1230 905 35 615 Ex. 5 Comp.
M 1.0 1225 890 65 620 Ex. 6 Comp.
N 1.2 1199 905 30 Ex. 7 Comp.
N 1.1 1273 900 35 Ex. 8 Comp.
N 1.3 1255 885 60 Ex. 9 Comp.
N 1.4 1230 895 50 Ex. 10 Comp.
N 1.1 1225 905 55 Ex. 11 Comp.
N 1.2 1240 910 45 Ex. 12 Comp.
N 0.9 1245 895 40 Ex. 13 Comp.
N 1.5 1238 885 60 Ex. 14 Table 2-2 Step (4) Step (5) Step (6) Annealing Starting Tempering Temper Heating Annealing Tempering No. soaking Rapid cooling temperature oftemperatur rolling rate time time temperature rate ( C/s) rapid cooling e reduction ( C/s) (s) (s) ( C) ( C) ( C) rate (%)
It should be noted that the chemical compositions of the ultra-high-strength dual-phase steel in Examples 1-7 and the related process parameters all met the control requirements of the design specification according to the present disclosure. The chemical compositions of the steels in Comparative Examples 1-6 all included parameters that failed to meet the requirements of the design according to the present disclosure. Although the chemical composition of steel grade N in Comparative Examples 7-14 met the requirements of the design according to the present disclosure, the related process parameters all included parameters that failed to meet the requirements of the design according to the present disclosure.
Tables 2-1 and 2-2 list the specific process parameters for the ultra-high-strength dual-phase steels in Examples 1-7 and the steels in Comparative Examples 1-14.
Table 2-1 Step (1) Step (2) Step (3) Drawing speed Finishing Steel Soaking Cooling Coiling Cold rolling No. in continuous rolling grade temperature rate temperature reduction casting temperature ( C) ( C/s) ( C) rate (%) (m/min) ( C) Ex. 1 A 1.0 1250 895 25 605 Ex. 2 B 1.2 1245 880 30 625 Ex. 3 c 1.5 1220 890 45 645 Ex. 4 D 0.9 1234 905 50 625 Ex. 5 E 1.1 1224 910 35 615 Ex. 6 F 1.3 1240 890 60 600 Ex. 7 G 1.0 1250 885 65 630 Comp. H 1.2 1230 895 40 615 Ex. 1 Comp.
I 0.9 1228 915 55 620 Ex. 2 Comp.
J 1.4 1250 890 70 625 Ex. 3 Comp.
K 1.3 1255 900 45 615 Ex. 4 Comp.
L 0.6 1230 905 35 615 Ex. 5 Comp.
M 1.0 1225 890 65 620 Ex. 6 Comp.
N 1.2 1199 905 30 Ex. 7 Comp.
N 1.1 1273 900 35 Ex. 8 Comp.
N 1.3 1255 885 60 Ex. 9 Comp.
N 1.4 1230 895 50 Ex. 10 Comp.
N 1.1 1225 905 55 Ex. 11 Comp.
N 1.2 1240 910 45 Ex. 12 Comp.
N 0.9 1245 895 40 Ex. 13 Comp.
N 1.5 1238 885 60 Ex. 14 Table 2-2 Step (4) Step (5) Step (6) Annealing Starting Tempering Temper Heating Annealing Tempering No. soaking Rapid cooling temperature oftemperatur rolling rate time time temperature rate ( C/s) rapid cooling e reduction ( C/s) (s) (s) ( C) ( C) ( C) rate (%)
11 Ex. 1 5 825 120 45 705 280 200 0.2 Ex. 2 8 820 75 35 670 290 300 0.1 Ex. 3 10 845 180 80 690 265 210 0.1 Ex. 4 9 805 100 55 700 305 250 0.3 Ex. 5 4 810 45 48 670 285 120 0.2 Ex. 6 3 828 55 56 680 290 300 0.3 Ex. 7 6 824 85 55 725 266 125 0.2 Comp. Ex. 1 6 832 150 48 695 286 330 0.1 Comp. Ex. 2 5 840 90 65 675 315 205 0.3 Comp. Ex. 3 8 800 105 70 705 307 190 0.1 Comp. Ex. 4 9 818 180 62 720 274 180 0.3 Comp. Ex. 5 5 835 60 58 715 264 240 0.2 Comp. Ex. 6 4 812 75 80 670 292 225 0.1 Comp. Ex. 7 8 834 120 55 720 305 325 0.2 Comp. Ex. 8 7 826 135 38 700 269 290 0.1 Comp. Ex. 9 6 819 95 80 680 296 175 0.2 Comp. Ex. 10 10 830 105 55 685 288 205 0.1 Comp. Ex. 11 3 794 145 62 695 308 380 0.1 Comp. Ex. 12 8 865 65 56 705 275 280 0.3 Comp. Ex. 13 5 807 175 48 710 340 300 0.2 Comp. Ex. 14 6 814 125 54 690 230 290 0.1 A variety of performance tests were performed on the ultra-high-strength dual-phase steels in Examples 1-7 and the steels in Comparative Examples 1-14. The test results obtained are listed in Table 3.
As to the performance test method, GB/T 13239-2006 Metallic Materials -Tensile Testing at Low Temperature was referred to. A standard sample was prepared, and subjected to static stretching on a tensile testing machine to obtain a corresponding stress-strain curve. After data processing, the parameters of yield strength, tensile strength and elongation after fracture were obtained finally.
Method for measurement of hydrogen content: The sample was heated to a certain temperature, and a hydrogen analyzer was used to measure the concentration of hydrogen released along with the change (rise) of the temperature, thereby judging the initial hydrogen content in the steel.
As to the performance test method, GB/T 13239-2006 Metallic Materials -Tensile Testing at Low Temperature was referred to. A standard sample was prepared, and subjected to static stretching on a tensile testing machine to obtain a corresponding stress-strain curve. After data processing, the parameters of yield strength, tensile strength and elongation after fracture were obtained finally.
Method for measurement of hydrogen content: The sample was heated to a certain temperature, and a hydrogen analyzer was used to measure the concentration of hydrogen released along with the change (rise) of the temperature, thereby judging the initial hydrogen content in the steel.
12 Table 3 lists the performance test results for the ultra-high-strength dual-phase steels in Examples 1-7 and the steels in Comparative Examples 1-14.
Table 3 Initial Yield Tensile Elongation hydrogen Stress level Stress level Stress level No. strength strength after fracture content 0.6*TS 0.8*TS
1.2*TS
(MPa) (MPa) (%) (PPm) Ex. 1 932 1329 9.7 5 0 0 Ex. 2 955 1338 9.2 7 0 0 Ex. 3 961 1340 8.5 3 0 0 Ex. 4 987 1364 7.6 6 0 0 Ex. 5 1004 1385 6.8 4 0 0 Ex. 6 1046 1407 6.2 1 0 0 Ex. 7 1065 1421 5.7 2 0 0 Comp. Ex. 1 877 1285 11.3 3 0 0 Comp. Ex. 2 1126 1469 4.3 9 0 X
X
Comp. Ex. 3 854 1276 11.5 2 0 0 Comp. Ex. 4 1134 1480 3.9 8 0 X
X
Comp. Ex. 5 865 1291 11.8 3 0 0 Comp. Ex. 6 839 1274 12.2 7 0 0 Comp. Ex. 7 852 1258 12.8 4 0 0 Comp. Ex. 8 1100 1446 4.7 8 0 0 X
Comp. Ex. 9 1098 1439 4.4 2 0 0 X
Comp. Ex. 10 870 1268 12.6 6 0 0 Comp. Ex. 11 886 1275 11.3 5 0 0 Comp. Ex. 12 1133 1485 3.9 8 0 0 X
Comp. Ex. 13 865 1286 11.6 6 0 0 Comp. Ex. 14 1108 1455 5.1 3 0 0 X
Note: The results of soaking the steel plates in 1 mol/L hydrochloric acid for 300 hours under a certain internal stress level: 0 represents no cracking, X represents cracking.
As it can be seen from Table 3, high-strength steels having a strength of at least 1300 MPa can be manufactured according to the present disclosure. Each Example according to the present disclosure has a yield strength of? 900 MPa, a tensile strength of? 1300 MPa, an elongation after
Table 3 Initial Yield Tensile Elongation hydrogen Stress level Stress level Stress level No. strength strength after fracture content 0.6*TS 0.8*TS
1.2*TS
(MPa) (MPa) (%) (PPm) Ex. 1 932 1329 9.7 5 0 0 Ex. 2 955 1338 9.2 7 0 0 Ex. 3 961 1340 8.5 3 0 0 Ex. 4 987 1364 7.6 6 0 0 Ex. 5 1004 1385 6.8 4 0 0 Ex. 6 1046 1407 6.2 1 0 0 Ex. 7 1065 1421 5.7 2 0 0 Comp. Ex. 1 877 1285 11.3 3 0 0 Comp. Ex. 2 1126 1469 4.3 9 0 X
X
Comp. Ex. 3 854 1276 11.5 2 0 0 Comp. Ex. 4 1134 1480 3.9 8 0 X
X
Comp. Ex. 5 865 1291 11.8 3 0 0 Comp. Ex. 6 839 1274 12.2 7 0 0 Comp. Ex. 7 852 1258 12.8 4 0 0 Comp. Ex. 8 1100 1446 4.7 8 0 0 X
Comp. Ex. 9 1098 1439 4.4 2 0 0 X
Comp. Ex. 10 870 1268 12.6 6 0 0 Comp. Ex. 11 886 1275 11.3 5 0 0 Comp. Ex. 12 1133 1485 3.9 8 0 0 X
Comp. Ex. 13 865 1286 11.6 6 0 0 Comp. Ex. 14 1108 1455 5.1 3 0 0 X
Note: The results of soaking the steel plates in 1 mol/L hydrochloric acid for 300 hours under a certain internal stress level: 0 represents no cracking, X represents cracking.
As it can be seen from Table 3, high-strength steels having a strength of at least 1300 MPa can be manufactured according to the present disclosure. Each Example according to the present disclosure has a yield strength of? 900 MPa, a tensile strength of? 1300 MPa, an elongation after
13 fracture of? 5%, and an initial hydrogen content of <10 ppm. The ultra-high-strength dual-phase steel in each Example has an ultra-high strength and a delayed cracking performance that is significantly better than that of a comparative steel grade of the same level.
No delayed cracking occurs when the steel plate is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength. The ultra-high-strength dual-phase steel in each Example has excellent performances. It is suitable for manufacture of automotive safety structural parts, and it is highly valuable and promising for popularization and application.
It's to be noted that the prior art portions in the protection scope of the present disclosure are not limited to the examples set forth in the present application file. All the prior art contents not contradictory to the technical solution of the present disclosure, including but not limited to prior patent literature, prior publications, prior public uses and the like, may all be incorporated into the protection scope of the present disclosure. In addition, the ways in which the various technical features of the present disclosure are combined are not limited to the ways recited in the claims of the present disclosure or the ways described in the specific examples. All the technical features recited in the present disclosure may be combined or integrated freely in any manner, unless contradictions are resulted.
It should also be noted that the Examples set forth above are only specific examples according to the present disclosure. Obviously, the present disclosure is not limited to the above Examples. Similar variations or modifications made thereto can be directly derived or easily contemplated from the present disclosure by those skilled in the art. They all fall in the protection scope of the present disclosure.
No delayed cracking occurs when the steel plate is soaked in 1 mol/L hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength. The ultra-high-strength dual-phase steel in each Example has excellent performances. It is suitable for manufacture of automotive safety structural parts, and it is highly valuable and promising for popularization and application.
It's to be noted that the prior art portions in the protection scope of the present disclosure are not limited to the examples set forth in the present application file. All the prior art contents not contradictory to the technical solution of the present disclosure, including but not limited to prior patent literature, prior publications, prior public uses and the like, may all be incorporated into the protection scope of the present disclosure. In addition, the ways in which the various technical features of the present disclosure are combined are not limited to the ways recited in the claims of the present disclosure or the ways described in the specific examples. All the technical features recited in the present disclosure may be combined or integrated freely in any manner, unless contradictions are resulted.
It should also be noted that the Examples set forth above are only specific examples according to the present disclosure. Obviously, the present disclosure is not limited to the above Examples. Similar variations or modifications made thereto can be directly derived or easily contemplated from the present disclosure by those skilled in the art. They all fall in the protection scope of the present disclosure.
14
Claims (15)
1. An ultra-high-strength dual-phase steel, wherein the ultra-high-strength dual-phase steel has a matrix structure of ferrite + martensite, wherein the ferrite and the martensite are distributed evenly like islands, and wherein the ultra-high-strength dual-phase steel comprises the following chemical elements in mass percentages, in addition to Fe:
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%4003%.
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%4003%.
2. The ultra-high-strength dual-phase steel according to claim 1, wherein the chemical elements have the following mass percentages:
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%-0.003%, and a balance of Fe and other unavoidable impurities.
C: 0.12-0.2%, Si: 0.5-1.0%, Mn: 2.5-3.0%, Al: 0.02-0.05%, Nb: 0.02-0.05%, Ti:
0.02-0.05%, B: 0.001%-0.003%, and a balance of Fe and other unavoidable impurities.
3. The ultra-high-strength dual-phase steel according to claim 2, wherein the unavoidable impurities include elements P, S and N, and contents thereof are controlled to be at least one of the following: P <0.01%, S<0.002%, N<0.004%.
4. The ultra-high-strength dual-phase steel according to any one of claims 1-3, wherein the mass percentages of the chemical elements satisfy at least one of:
C: 0.14-0.18%, Mn: 2.5-2.8%.
C: 0.14-0.18%, Mn: 2.5-2.8%.
5. The ultra-high-strength dual-phase steel according to any one of claims 1-3, wherein the martensite has a phase proportion of >90%.
6. The ultra-high-strength dual-phase steel according to any one of claims 1-3, wherein the martensite comprises coherently distributed E carbides.
7. The ultra-high-strength dual-phase steel according to any one of claims 1-3, wherein the ultra-high-strength dual-phase steel has performances that meet at least one of the following:
yield strength 2900 MPa, tensile strength 21300 MPa, elongation after fracture 25%, initial hydrogen content <10 ppm; no delayed cracking when soaked in 1 mol/L
hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength.
yield strength 2900 MPa, tensile strength 21300 MPa, elongation after fracture 25%, initial hydrogen content <10 ppm; no delayed cracking when soaked in 1 mol/L
hydrochloric acid for 300 hours under a pre-stress of greater than or equal to the tensile strength.
8. The ultra-high-strength dual-phase steel according to any one of claims 1-3, wherein the ultra-high-strength dual-phase steel has performances that meet at least one of the following:
yield strength 2930 MPa, tensile strength 21320 MPa, elongation after fracture 25.5%, initial hydrogen content <7 ppm; no delayed cracking when soaked in 1 mol/L
hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength.
yield strength 2930 MPa, tensile strength 21320 MPa, elongation after fracture 25.5%, initial hydrogen content <7 ppm; no delayed cracking when soaked in 1 mol/L
hydrochloric acid for 300 hours under a pre-stress of greater than or equal to 1.2 times of the tensile strength.
9. The ultra-high-strength dual-phase steel according to any one of claims 1-3, wherein the ultra-high-strength dual-phase steel has a yield ratio of 0.70-0.75.
10. A manufacturing method for the ultra-high-strength dual-phase steel according to any one of claims 1-9, wherein the method comprises steps of:
(1) Smelting and continuous casting;
(2) Hot rolling;
(3) Cold rolling;
(4) Annealing: heating to an annealing soaking temperature of 800-850 C at a heating rate of 3-10 C/s, the annealing time being 40-200 s; and then rapidly cooling at a rate of 30-80 C/s, a starting temperature of the rapid cooling being 670-730 C;
(5) Tempering: tempering temperature: 260-320 C; tempering time: 100-400 s;
(6) Temper rolling;
(7) Electrogalvanizing.
(1) Smelting and continuous casting;
(2) Hot rolling;
(3) Cold rolling;
(4) Annealing: heating to an annealing soaking temperature of 800-850 C at a heating rate of 3-10 C/s, the annealing time being 40-200 s; and then rapidly cooling at a rate of 30-80 C/s, a starting temperature of the rapid cooling being 670-730 C;
(5) Tempering: tempering temperature: 260-320 C; tempering time: 100-400 s;
(6) Temper rolling;
(7) Electrogalvanizing.
11. The manufacturing method according to claim 10, wherein in step (1), a drawing speed in the continuous casting is controlled at 0.9-1.5 m/min during the continuous casting process.
12. The manufacturing method according to claim 10, wherein in step (2), a cast slab is controlled to be soaked at a temperature of 1220-1260 C; then rolled with a finishing rolling temperature being controlled at 880-920 C; then cooled at a rate of 20-70 C/s after the rolling;
then coiled at a coiling temperature of 600-650 C; and then subjected to heat preservation treatment after the coiling.
then coiled at a coiling temperature of 600-650 C; and then subjected to heat preservation treatment after the coiling.
13. The manufacturing method according to claim 10, wherein in step (3), a cold rolling reduction rate is controlled at 45-65%.
14. The manufacturing method according to claim 10, wherein in step (6), a temper rolling reduction rate is controlled at <0.3%; and/or in step (7), double-side electrogalvanizing is performed with a plating layer weight of 10-100 g/m2 on each side.
15. The manufacturing method according to claim 10, wherein in step (2), a cast slab is controlled to be soaked at a temperature of 1220-1250 C, and a coiling temperature is 605-645 C; in step (4), the annealing soaking temperature is 805-845 C; in step (5), the tempering temperature is 260-310 C, and the tempering time is preferably 100-300 s.
Applications Claiming Priority (3)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| CN202010459510.0 | 2020-05-27 | ||
| CN202010459510.0A CN113737087B (en) | 2020-05-27 | 2020-05-27 | Ultrahigh-strength dual-phase steel and manufacturing method thereof |
| PCT/CN2021/095807 WO2021238916A1 (en) | 2020-05-27 | 2021-05-25 | Ultrahigh-strength dual-phase steel and manufacturing method therefor |
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| CA3180467A1 true CA3180467A1 (en) | 2021-12-02 |
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| CA3180467A Pending CA3180467A1 (en) | 2020-05-27 | 2021-05-25 | Ultrahigh-strength dual-phase steel and manufacturing method therefor |
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| US (1) | US20230227930A1 (en) |
| EP (1) | EP4159886A4 (en) |
| JP (1) | JP2023527389A (en) |
| CN (1) | CN113737087B (en) |
| CA (1) | CA3180467A1 (en) |
| WO (1) | WO2021238916A1 (en) |
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| CN115216708B (en) * | 2022-06-17 | 2023-08-22 | 攀钢集团攀枝花钢铁研究院有限公司 | Low-cost high-yield-ratio 1200MPa grade cold-rolled dual-phase steel and preparation method thereof |
| CN117305683A (en) * | 2022-06-22 | 2023-12-29 | 宝山钢铁股份有限公司 | Cold-rolled steel plate with grade of 1300MPa or more and manufacturing method thereof |
| CN117660831A (en) * | 2022-08-23 | 2024-03-08 | 宝山钢铁股份有限公司 | Dual-phase steel and manufacturing method thereof |
| CN117660830A (en) * | 2022-08-23 | 2024-03-08 | 宝山钢铁股份有限公司 | 100 kg-level cold-rolled low-alloy annealed dual-phase steel and manufacturing method thereof |
| CN117660846A (en) * | 2022-08-23 | 2024-03-08 | 宝山钢铁股份有限公司 | 120 kg-level cold-rolled low-alloy annealed dual-phase steel and manufacturing method thereof |
| CN117089761B (en) * | 2023-05-30 | 2024-06-04 | 宝山钢铁股份有限公司 | Variable-strength dual-phase steel plate with normalized components and flexible manufacturing method thereof |
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| JPH11270531A (en) * | 1998-03-19 | 1999-10-05 | Nippon Steel Corp | High strength bolt having good delayed fracture characteristic and manufacture thereof |
| CN100392135C (en) * | 2005-06-30 | 2008-06-04 | 宝山钢铁股份有限公司 | Ultra-high strength strip steel and its production process |
| CN100430505C (en) * | 2005-09-29 | 2008-11-05 | 宝山钢铁股份有限公司 | Superhigh-strength cold rolling band steel with anti-tensile strength above 880Mpa and its production |
| EP1951519A4 (en) * | 2005-10-24 | 2008-12-31 | Exxonmobil Upstream Res Co | High strength dual phase steel with low yield ratio, high toughness and superior weldability |
| JP5668337B2 (en) * | 2010-06-30 | 2015-02-12 | Ｊｆｅスチール株式会社 | Ultra-high-strength cold-rolled steel sheet excellent in ductility and delayed fracture resistance and method for producing the same |
| EP2803747B1 (en) * | 2012-01-13 | 2019-03-27 | Nippon Steel & Sumitomo Metal Corporation | Cold-rolled steel sheet and method for producing cold-rolled steel sheet |
| JP6048382B2 (en) * | 2013-12-09 | 2016-12-21 | Ｊｆｅスチール株式会社 | High-strength cold-rolled steel sheet and manufacturing method thereof |
| CN107208222B (en) * | 2015-01-28 | 2018-11-27 | 杰富意钢铁株式会社 | High strength cold rolled steel plate, high-intensitive coated steel sheet and its manufacturing method |
| WO2018220412A1 (en) * | 2017-06-01 | 2018-12-06 | Arcelormittal | Method for producing high-strength steel parts with improved ductility, and parts obtained by said method |
| CN109207867A (en) * | 2017-06-29 | 2019-01-15 | 宝山钢铁股份有限公司 | A kind of cold rolled annealed dual phase steel, steel plate and its manufacturing method |
| CN109207841B (en) * | 2017-06-30 | 2021-06-15 | 宝山钢铁股份有限公司 | Low-cost high-formability 1180 MPa-grade cold-rolled annealed dual-phase steel plate and manufacturing method thereof |
| CN111247264A (en) * | 2017-10-20 | 2020-06-05 | 杰富意钢铁株式会社 | High-strength steel sheet and method for producing same |
| CN108486494B (en) | 2018-06-05 | 2019-06-21 | 西北工业大学 | The production method of vanadium micro-alloying 1300MPa grade high-strength hot rolled steel plate and cold-rolled biphase steel plate |
| WO2020045220A1 (en) * | 2018-08-31 | 2020-03-05 | Ｊｆｅスチール株式会社 | High-strength steel plate and method for producing same |
| CN109280857B (en) * | 2018-09-25 | 2021-03-05 | 邯郸钢铁集团有限责任公司 | 1200 MPa-grade ultrafast cold-rolled dual-phase steel plate and preparation method thereof |
| CN109628846B (en) | 2018-12-20 | 2020-08-04 | 唐山钢铁集团有限责任公司 | 1300 MPa-grade ultrahigh-strength cold-rolled steel plate for automobiles and production method thereof |
| CN109504930B (en) | 2018-12-20 | 2020-10-02 | 唐山钢铁集团有限责任公司 | Hot-dip galvanized steel sheet with tensile strength of more than 1300MPa and production method thereof |
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- 2021-05-25 JP JP2022572701A patent/JP2023527389A/en active Pending
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| Publication number | Publication date |
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| JP2023527389A (en) | 2023-06-28 |
| EP4159886A1 (en) | 2023-04-05 |
| WO2021238916A1 (en) | 2021-12-02 |
| EP4159886A4 (en) | 2024-04-17 |
| US20230227930A1 (en) | 2023-07-20 |
| CN113737087A (en) | 2021-12-03 |
| CN113737087B (en) | 2022-07-19 |
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