CN113166893B - High-strength steel material having excellent durability and method for producing same - Google Patents
High-strength steel material having excellent durability and method for producing same Download PDFInfo
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- CN113166893B CN113166893B CN201980078073.3A CN201980078073A CN113166893B CN 113166893 B CN113166893 B CN 113166893B CN 201980078073 A CN201980078073 A CN 201980078073A CN 113166893 B CN113166893 B CN 113166893B
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 149
- 239000010959 steel Substances 0.000 title claims abstract description 149
- 239000000463 material Substances 0.000 title claims abstract description 58
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 27
- 238000000034 method Methods 0.000 claims abstract description 34
- 238000001816 cooling Methods 0.000 claims description 88
- 238000005098 hot rolling Methods 0.000 claims description 42
- 229910000859 α-Fe Inorganic materials 0.000 claims description 36
- 230000014509 gene expression Effects 0.000 claims description 34
- 239000011572 manganese Substances 0.000 claims description 33
- 239000010955 niobium Substances 0.000 claims description 32
- 239000010936 titanium Substances 0.000 claims description 32
- 239000011651 chromium Substances 0.000 claims description 28
- 239000013078 crystal Substances 0.000 claims description 28
- 229910000734 martensite Inorganic materials 0.000 claims description 23
- 229910001563 bainite Inorganic materials 0.000 claims description 21
- IJGRMHOSHXDMSA-UHFFFAOYSA-N Atomic nitrogen Chemical compound N#N IJGRMHOSHXDMSA-UHFFFAOYSA-N 0.000 claims description 14
- 238000005096 rolling process Methods 0.000 claims description 12
- XAGFODPZIPBFFR-UHFFFAOYSA-N aluminium Chemical compound [Al] XAGFODPZIPBFFR-UHFFFAOYSA-N 0.000 claims description 10
- 229910052782 aluminium Inorganic materials 0.000 claims description 10
- 238000010438 heat treatment Methods 0.000 claims description 10
- 239000012535 impurity Substances 0.000 claims description 10
- 229910052719 titanium Inorganic materials 0.000 claims description 9
- 238000004804 winding Methods 0.000 claims description 9
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 claims description 8
- 229910052799 carbon Inorganic materials 0.000 claims description 8
- RTAQQCXQSZGOHL-UHFFFAOYSA-N Titanium Chemical compound [Ti] RTAQQCXQSZGOHL-UHFFFAOYSA-N 0.000 claims description 7
- 229910001566 austenite Inorganic materials 0.000 claims description 7
- 229910052758 niobium Inorganic materials 0.000 claims description 7
- GUCVJGMIXFAOAE-UHFFFAOYSA-N niobium atom Chemical compound [Nb] GUCVJGMIXFAOAE-UHFFFAOYSA-N 0.000 claims description 7
- 229910052757 nitrogen Inorganic materials 0.000 claims description 7
- 229910052710 silicon Inorganic materials 0.000 claims description 7
- 229910052717 sulfur Inorganic materials 0.000 claims description 7
- VYZAMTAEIAYCRO-UHFFFAOYSA-N Chromium Chemical compound [Cr] VYZAMTAEIAYCRO-UHFFFAOYSA-N 0.000 claims description 6
- PWHULOQIROXLJO-UHFFFAOYSA-N Manganese Chemical compound [Mn] PWHULOQIROXLJO-UHFFFAOYSA-N 0.000 claims description 6
- OAICVXFJPJFONN-UHFFFAOYSA-N Phosphorus Chemical compound [P] OAICVXFJPJFONN-UHFFFAOYSA-N 0.000 claims description 6
- XUIMIQQOPSSXEZ-UHFFFAOYSA-N Silicon Chemical compound [Si] XUIMIQQOPSSXEZ-UHFFFAOYSA-N 0.000 claims description 6
- NINIDFKCEFEMDL-UHFFFAOYSA-N Sulfur Chemical compound [S] NINIDFKCEFEMDL-UHFFFAOYSA-N 0.000 claims description 6
- 229910052804 chromium Inorganic materials 0.000 claims description 6
- 229910052748 manganese Inorganic materials 0.000 claims description 6
- 229910052698 phosphorus Inorganic materials 0.000 claims description 6
- 239000011574 phosphorus Substances 0.000 claims description 6
- 239000010703 silicon Substances 0.000 claims description 6
- 239000011593 sulfur Substances 0.000 claims description 6
- 238000005246 galvanizing Methods 0.000 claims description 4
- 239000011777 magnesium Substances 0.000 claims description 4
- 238000007747 plating Methods 0.000 claims description 3
- FYYHWMGAXLPEAU-UHFFFAOYSA-N Magnesium Chemical compound [Mg] FYYHWMGAXLPEAU-UHFFFAOYSA-N 0.000 claims description 2
- 229910052749 magnesium Inorganic materials 0.000 claims description 2
- 238000005554 pickling Methods 0.000 claims 2
- 230000000052 comparative effect Effects 0.000 description 35
- 230000000694 effects Effects 0.000 description 20
- 239000000203 mixture Substances 0.000 description 18
- 229910045601 alloy Inorganic materials 0.000 description 15
- 239000000956 alloy Substances 0.000 description 15
- 229910001562 pearlite Inorganic materials 0.000 description 13
- 230000015572 biosynthetic process Effects 0.000 description 11
- 238000005728 strengthening Methods 0.000 description 11
- 230000009466 transformation Effects 0.000 description 11
- 239000002244 precipitate Substances 0.000 description 9
- 238000001953 recrystallisation Methods 0.000 description 8
- 150000001247 metal acetylides Chemical class 0.000 description 7
- 239000006104 solid solution Substances 0.000 description 6
- 230000006866 deterioration Effects 0.000 description 5
- 238000009661 fatigue test Methods 0.000 description 5
- 239000011159 matrix material Substances 0.000 description 5
- 238000001556 precipitation Methods 0.000 description 5
- 238000005204 segregation Methods 0.000 description 5
- 238000012360 testing method Methods 0.000 description 5
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 4
- 239000002131 composite material Substances 0.000 description 4
- 229910001568 polygonal ferrite Inorganic materials 0.000 description 4
- 239000002344 surface layer Substances 0.000 description 4
- 238000005452 bending Methods 0.000 description 3
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 3
- 238000000465 moulding Methods 0.000 description 3
- 230000000644 propagated effect Effects 0.000 description 3
- 238000004080 punching Methods 0.000 description 3
- 239000011701 zinc Substances 0.000 description 3
- HCHKCACWOHOZIP-UHFFFAOYSA-N Zinc Chemical compound [Zn] HCHKCACWOHOZIP-UHFFFAOYSA-N 0.000 description 2
- 238000009749 continuous casting Methods 0.000 description 2
- 230000007423 decrease Effects 0.000 description 2
- 230000001934 delay Effects 0.000 description 2
- 230000003111 delayed effect Effects 0.000 description 2
- 238000005530 etching Methods 0.000 description 2
- 230000000704 physical effect Effects 0.000 description 2
- 238000009628 steelmaking Methods 0.000 description 2
- 229910052725 zinc Inorganic materials 0.000 description 2
- 241000219307 Atriplex rosea Species 0.000 description 1
- 229910001297 Zn alloy Inorganic materials 0.000 description 1
- 230000002159 abnormal effect Effects 0.000 description 1
- 238000006243 chemical reaction Methods 0.000 description 1
- 238000005520 cutting process Methods 0.000 description 1
- 230000007547 defect Effects 0.000 description 1
- 238000009826 distribution Methods 0.000 description 1
- 239000012467 final product Substances 0.000 description 1
- 230000002401 inhibitory effect Effects 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 230000003287 optical effect Effects 0.000 description 1
- 239000000047 product Substances 0.000 description 1
- 230000001737 promoting effect Effects 0.000 description 1
- 239000002994 raw material Substances 0.000 description 1
- 238000007670 refining Methods 0.000 description 1
- 230000003014 reinforcing effect Effects 0.000 description 1
- 238000010583 slow cooling Methods 0.000 description 1
- 239000013585 weight reducing agent Substances 0.000 description 1
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- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
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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
- 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
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
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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
- 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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- 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
- C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
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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
- C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
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- C21D6/00—Heat treatment of ferrous alloys
- C21D6/002—Heat treatment of ferrous alloys containing Cr
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- C21D6/00—Heat treatment of ferrous alloys
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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
- C21D8/02—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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- 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
- C21D8/0231—Warm rolling
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- C22C18/04—Alloys based on zinc with aluminium as the next major constituent
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- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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- 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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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
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- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
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- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
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- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
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- C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
- C23C2/0224—Two or more thermal pretreatments
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- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
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- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
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- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
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- C21D2211/00—Microstructure comprising significant phases
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Abstract
The present invention relates to a steel material used for a frame member of a chassis member of a commercial vehicle, a wheel disc, or the like, and more particularly, to a high-strength steel material having excellent durability and a method for manufacturing the same.
Description
Technical Field
The present invention relates to a high-strength steel material having excellent durability and a method for producing the same.
Background
Conventionally, high-strength steel sheets having a thickness of 5mm or more and a yield strength of 450 to 600MPa have been used for frame members and wheel discs of chassis members of commercial vehicles to ensure high rigidity according to vehicle characteristics, but in recent years, high-strength steel materials having a tensile strength of 650MPa or more have been used for weight reduction and high strength of automobiles.
When a high-strength steel material is used to manufacture a part, there is a disadvantage that a durability life of a final product (part) is shortened because fine cracks are formed at a sheared portion of a steel sheet through a step of press forming a sheet material subjected to shear forming and punch forming.
As a method for solving such a problem, patent document 1 discloses a method in which a ferrite phase is used as a matrix structure and fine precipitates are formed by rolling the ferrite phase at a high temperature after a conventional austenite hot rolling. Patent document 2 proposes a technique of cooling the winding temperature to a temperature at which a bainite phase is formed as a matrix structure, and then winding the bainite phase so that a coarse pearlite structure is not formed. Patent document 3 discloses a technique of performing rolling at a reduction ratio of 40% or more in a non-recrystallized region in a hot rolling process by using titanium (Ti), niobium (Nb), or the like to refine austenite grains.
In order to manufacture high-strength steel, alloy components such as Si, mn, al, mo, cr, etc. are mainly used, and in this case, the strength of the hot-rolled steel sheet is effectively improved, but when a large amount of alloy components is added, a part of the components segregate (segregation) in the steel or cause unevenness in a microstructure, so that shear formability is deteriorated, and fine cracks generated at a shear surface are easily propagated in a fatigue environment, so that breakage of components occurs.
In particular, as the thickness of the steel material increases, the unevenness of the microstructure between the thickness surface layer portion and the central portion increases, so that the generation of cracks at the shear plane increases, and the propagation speed of the cracks becomes faster in a fatigue environment, resulting in deterioration of durability.
However, the above-mentioned techniques (patent documents 1 to 3) do not consider the fatigue characteristics of thick steel materials having high strength.
In addition, when a precipitation strengthening effect is obtained by refining crystal grains of a thick steel material using a precipitation forming element such as Ti, nb, or V, if coiling is performed at a high temperature of about 500 to 700 ℃ at which precipitates are easily formed, or if the cooling rate of a steel sheet is not controlled in a cooling process after hot rolling, coarse carbides are formed in the thickness center portion of the thick steel material, and the quality of a shear surface deteriorates. Also, the application of a large reduction of 40% to the unrecrystallized region during hot rolling results in deterioration of the shape quality of the rolled sheet and loads on equipment, thus presenting a problem of difficulty in practical use.
(patent document 1) Japanese laid-open patent publication No. 2002-322541
(patent document 2) Korean patent laid-open publication No. 10-1528084
(patent document 3) Japanese laid-open patent publication No. 1997-143570
Disclosure of Invention
Technical problem to be solved
An object of one aspect of the present invention is to provide a steel material that is thick and has a certain thickness, high strength, and excellent durability, and a method for manufacturing the same.
The technical problem of the present invention is not limited to the above. It is not difficult for those skilled in the art to understand the additional technical problems of the present invention from the contents of the entire specification of the present invention.
Technical scheme
One aspect of the present invention provides a high-strength steel material having excellent durability, the high-strength steel material comprising, in wt%: carbon (C): 0.05-0.15%, silicon (Si): 0.01-1.0%, manganese (Mn): 1.0-2.3%, aluminum (Al): 0.01-0.1%, chromium (Cr): 0.005-1.0%, phosphorus (P): 0.001-0.05%, sulfur (S): 0.001-0.01%, nitrogen (N): 0.001-0.01%, niobium (Nb): 0.005-0.07%, titanium (Ti): 0.005-0.11%, and the balance Fe and other unavoidable impurities, the sum of fractions of a ferrite phase and a bainite phase being 90% or more as a fine structure, the fraction of crystal grains having an aspect ratio (ratio of short side/long side) of 0.3 or less in a central portion (t/4 position to t/2 position in the thickness direction) being less than 50%, and the unit area (1 mm) in the central portion 2 ) The length of the grain boundary observed therein is 700mm or more.
Another aspect of the present invention provides a method for manufacturing a high-strength steel material having excellent durability, including the steps of: at 1200-heating a steel blank having the above alloy composition at a temperature in the range of-1350 ℃; hot rolling the heated slab to manufacture a hot rolled steel sheet; cooling the hot rolled steel plate to a temperature range of 400-500 ℃ and then rolling (CT); and air-cooling to a temperature range of normal temperature to 200 ℃ after the rolling, wherein the hot rolling is performed under the condition that the following [ relational expression 1] is satisfied]Is subjected to finish hot rolling at a temperature (FDT (DEG C)). The cooling is subjected to primary cooling to satisfy [ relational expression 2] and secondary cooling]Cooling Rate (CR) 1 ) Is carried out so as to satisfy the following [ relational expression 3]]Cooling Rate (CR) 2 ) The process is carried out.
[ relational expression 1]
FDT (hot rolling termination temperature (DEG C)) of Tn-50 or less and Tn or less
Tn =730+92 × [ C ] +70 × [ Mn ] +45 × [ Cr ] +650 × [ Nb ] +410 × [ Ti ] -80 × [ Si ] -1.4 × (t-5) (where each element represents its weight percentage and t represents the thickness (mm) of the final hot-rolled steel sheet),
[ relational expression 2]
CR 1 ≥196-300×[C]+4.5×[Si]-71.8×[Mn]-59.6×[Cr]+187×[Ti]+852×[Nb](wherein, each element represents the weight percentage content),
[ relational expression 3]
CR Minimum size ≤CR 2 ≤CR Maximum of
(CR Maximum of ＝76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[Nb]，CR Minimum size of ＝27.4-45.3×[C]+5.28×[Si]-11×[Mn]-7.33×[Cr]+42.3×[Ti]+82×[Nb]Each element represents a weight percentage thereof).
Advantageous effects
According to the present invention, since the thick steel material has high strength and excellent quality of a cross section at the time of forming, an excellent ratio of the fatigue limit to the yield strength of the steel material can be secured after forming.
The steel material of the present invention has an effect that it can be suitably used for a skeleton member, a wheel disc, and the like of a chassis member of an automobile.
Drawings
Fig. 1 is a graph showing the ratio of fatigue strength to yield strength according to thickness of inventive steel and comparative steel according to an embodiment of the present invention.
Best mode for carrying out the invention
The present inventors have conducted intensive studies to solve the problem of the decrease in durability of conventional thick steel materials for automobiles during molding.
In particular, the present inventors examined the change in the crack distribution and durability of the conventional thick steel material according to the composition and the shear plane after forming of the microstructure, and confirmed that the durability characteristics change according to the control of the shape of the crystal grains in the thickness center portion of the steel material.
Thus, the present inventors have confirmed that a steel material having high strength and excellent quality of a cross section at the time of forming and thus having desired durability can be provided, and have completed the present invention.
The present invention will be described in detail below.
The high-strength steel material excellent in durability according to one aspect of the present invention may include, in wt%: carbon (C): 0.05-0.15%, silicon (Si): 0.01-1.0%, manganese (Mn): 1.0-2.3%, aluminum (Al): 0.01-0.1%, chromium (Cr): 0.005-1.0%, phosphorus (P): 0.001-0.05%, sulfur (S): 0.001-0.01%, nitrogen (N): 0.001-0.01%, niobium (Nb): 0.005-0.07%, titanium (Ti): 0.005-0.11%.
Hereinafter, the reason why the alloy composition of the hot rolled steel sheet provided in the present invention is limited as described above will be described in detail.
In addition, unless otherwise specified, the contents of the respective elements in the present invention are based on weight, and the proportion of the structure is based on area.
Carbon (C): 0.05 to 0.15 percent
Carbon (C) is the most economical and effective element for strengthening steel, and when the amount of C added increases, the precipitation strengthening effect increases, or the fraction of bainite phase increases, thereby improving tensile strength. Further, as the thickness of the hot-rolled steel material becomes thicker, the cooling rate of the thickness center portion becomes slower in the cooling process after hot rolling, and therefore, when the content of C is large, coarse carbides or pearlite is easily formed.
In the present invention, when the content of C is less than 0.05%, it is difficult to sufficiently obtain the reinforcing effect of the steel, while when the content of C exceeds 0.15%, pearlite phase or coarse carbide is formed in the central portion of the thickness, and thus there is a problem that shear formability is deteriorated and durability is lowered.
Therefore, in the present invention, the C may be contained in an amount of 0.05 to 0.15%, and more preferably, 0.06 to 0.12%.
Silicon (Si): 0.01 to 1.0 percent
Silicon (Si) deoxidizes molten steel and has a solid solution strengthening effect, and delays the formation of coarse carbides, thus contributing to improved formability.
When the content of Si is less than 0.01%, the solid solution strengthening effect is small, and the effect of delaying the formation of carbides is also reduced, so that it is difficult to improve the formability. On the other hand, when the content of Si exceeds 1.0%, red scale due to Si is formed on the surface of the steel sheet during hot rolling, resulting in a problem that the surface quality of the steel sheet becomes very poor and ductility and weldability are lowered.
Therefore, in the present invention, the Si may be contained in an amount of 0.01 to 1.0%, and more preferably, 0.2 to 0.7%.
Manganese (Mn): 1.0 to 2.3 percent
As Si, manganese (Mn) is an element effective for solid solution strengthening of steel, and increases hardenability of steel, so that a bainite phase is easily formed in cooling after hot rolling.
When the Mn content is less than 1.0%, the above-described effects cannot be sufficiently obtained. On the other hand, when the Mn content exceeds 2.3%, hardenability is greatly increased, so martensite transformation is easily caused, and when a slab is cast in a continuous casting process, a segregation portion at the thickness center portion becomes very developed, and when cooled after hot rolling, a fine structure is unevenly formed in the thickness direction, so shear formability and durability become poor.
Therefore, in the present invention, 1.0 to 2.3% of the Mn may be contained, and more preferably, 1.1 to 2.0% of the Mn may be contained.
Aluminum (Al): 0.01 to 0.1 percent
Aluminum (Al) is an element mainly added for deoxidation, and when the content of Al is less than 0.01%, the addition effect cannot be sufficiently obtained. On the other hand, when the content of Al exceeds 0.1%, al is combined with nitrogen (N) in the steel to form AlN, so that corner cracks are easily generated in the slab at the time of continuous casting, and defects due to the formation of inclusions may be generated.
Therefore, the Al may be contained in an amount of 0.01 to 0.1% in the present invention.
In the present invention, aluminum represents acid-soluble aluminum (sol. Al).
Chromium (Cr): 0.005-1.0%
Chromium (Cr) solution-strengthens steel and delays transformation of a ferrite phase upon cooling, thereby serving as a contribution to bainite formation at a winding temperature. In order to obtain the above effects, it is preferable to contain Cr of 0.005% or more, but when the content of Cr exceeds 1.0%, ferrite transformation is excessively delayed to form a martensite phase, resulting in deterioration of elongation. Further, similarly to Mn, the segregation portion at the thickness center portion becomes very developed, and the fine structure in the thickness direction becomes uneven, so that the shear formability and durability become poor.
Therefore, in the present invention, the Cr may be contained in an amount of 0.005 to 1.0%, and more preferably, 0.3 to 0.9%.
Phosphorus (P): 0.001-0.05%
Phosphorus (P) is an element having both a solid-solution strengthening effect and a ferrite transformation promoting effect. In order to manufacture the content of P less than 0.001%, the manufacturing cost required is excessively high, and thus it is economically disadvantageous, and it is also difficult to secure the strength of the target level. When the content of P exceeds 0.05%, brittleness due to grain boundary segregation is generated, fine cracks are easily generated during molding, and shear moldability and durability are greatly reduced.
Therefore, the present invention may contain 0.001 to 0.05% of said P.
Sulfur (S): 0.001-0.01%
Sulfur (S) is an impurity present in steel, and when the content of S exceeds 0.01%, S bonds with Mn and the like to form a non-metallic inclusion, so that fine cracks are easily generated at the time of cutting of steel, and there is a problem that shear formability and durability are greatly reduced. In addition, in order to produce the steel product with the S content of less than 0.001%, an excessive amount of time is required for the steel making operation, and thus productivity is reduced.
Therefore, the present invention may contain 0.001 to 0.01% of said S.
Nitrogen (N): 0.001-0.01%
Like C, nitrogen (N) is a representative solid solution strengthening element, and the N bonds with Ti, al, and the like to form coarse precipitates. In general, N is superior in solid solution strengthening effect to carbon, but as the amount of N in steel increases, there is a problem that toughness of steel decreases. In view of this, it is preferable to contain 0.01% or less of the N, but in order to produce the steel with the N content of less than 0.001%, a large amount of time is required for the steel-making operation, and thus productivity is lowered.
Therefore, the present invention may contain 0.001 to 0.01% of said N.
Niobium (Nb): 0.005-0.07%
Niobium (Nb) is a precipitation strengthening element, precipitates Nb during hot rolling, and effectively improves the strength and impact toughness of steel by the effect of grain refinement due to recrystallization retardation. In order to sufficiently obtain the above effects, the Nb may be contained at 0.005% or more, and on the other hand, when the content of the Nb exceeds 0.07%, the formability and durability are deteriorated due to the formation of elongated crystal grains and the formation of coarse composite precipitates caused by excessive delay of recrystallization during hot rolling.
Therefore, in the present invention, 0.005 to 0.07% of the Nb may be contained, and more preferably, 0.01 to 0.06% of the Nb may be contained.
Titanium (Ti): 0.005-0.11%
As the Nb, titanium (Ti) is a representative precipitation strengthening element, and the Ti forms coarse TiN in steel by a strong affinity with N. The TiN has an effect of inhibiting grain growth during heating for hot rolling. In addition, ti remaining after the reaction with N is dissolved in the steel and combined with carbon to form TiC precipitates, which is advantageous for improving the strength of the steel.
In order to sufficiently obtain the above effects, it is necessary to contain 0.005% or more of Ti, but when the content of Ti exceeds 0.11%, there is a problem that collision resistance during molding is deteriorated due to generation of coarse TiN and coarsening of precipitates.
Therefore, in the present invention, the Ti may be contained in an amount of 0.005 to 0.11%, and more preferably, 0.01 to 0.1%.
The remainder of the composition of the invention is iron (Fe). However, since unnecessary impurities are inevitably mixed from the raw materials or the surrounding environment in a normal manufacturing process, the impurities cannot be excluded. These impurities are well known to the skilled person in the usual manufacturing process and therefore not all of them are specifically mentioned in this specification.
The microstructure of the steel material of the present invention having the alloy composition may be a composite structure of a ferrite phase and a bainite phase.
In this case, the sum of the area fractions of the ferrite phase and the bainite phase is preferably 90% or more, and the area fraction of the bainite phase may be 50% or more.
When the area fraction of the bainite phase is less than 50%, it is difficult to secure a desired strength, and when a coarse ferrite phase increases, it has an uneven fine structure, so that fine cracks are likely to be generated during shear deformation or punch deformation.
The ferrite phase refers to a polygonal ferrite phase as a high-temperature-region ferrite phase, and the bainite phase is a general term for all acicular ferrite and bainitic ferrite phases as low-temperature-region ferrite phases.
The balance structure other than the composite structure may contain a MA phase (mixed structure of martensite and austenite) and a martensite phase. At this time, the sum of the area fractions of the MA phase and the martensite phase may be included in the range of 1 to 10%, wherein the MA phase is preferably less than 3%.
When the sum of the fractions of the MA phase and the martensite phase exceeds 10%, the tensile strength increases, but the hardness value is higher than that of the surrounding structure phase (phase), and therefore, cracks are generated at the interface between the MA phase and the martensite phase at the time of shear deformation or punch deformation, resulting in deterioration of fatigue characteristics. In particular, the average size of the MA phase is on the order of 1/10 as compared with the martensite phase, but the area fraction is preferably limited to 3% or less because cracks tend to form at the interface of the phases, and the propagation rate increases when exposed to a fatigue environment, similarly to the martensite phase.
As described above, the effect of eliminating the structural unevenness can be obtained by minimizing the fractions of the coarse MA phase and the martensite phase in the matrix structure.
In addition, in addition to the above-mentioned structure, even if the steel material of the present invention contains a pearlite phase in an amount of 3% or less (including 0%), there is no great difficulty in securing desired physical properties.
In particular, the fraction of crystal grains having an aspect ratio (ratio of short side length (minor axis)/long side length (major axis)) of 0.3 or less in the central portion corresponding to the position t/4 to the position t/2 in the thickness direction, the unit area (1 mm) of the central portion in the steel material of the present invention is preferably less than 50% 2 ) The length of the grain boundary observed inside is preferably 700mm or more.
When the fraction of crystal grains having an aspect ratio of 0.3 or less in the central portion is 50% or more, cracks easily grow when cracks occur, and thus the durability is deteriorated. Further, when the length of the grain boundary in the central portion is less than 700mm, the strength of the central portion is reduced, and cracks are easily propagated, so that the durability may also be deteriorated.
The method for analyzing the aspect ratio of the crystal grains and the length of the grain boundary is not particularly limited, but analysis may be performed by Electron Back Scattered Diffraction (EBSD) as an example. Specifically, in the EBSD measurement results of the rolled section, the area per unit area (1 mm) was calculated from the crystal grains having a large-angle grain boundary of 15 ° or more 2 ) The length and aspect ratio of the grain boundary of (2) can be calculated from the ratio of the minor axis to the major axis of the grain size.
The steel material of the present invention having the alloy composition and the microstructure described above is a thick steel material having a thickness of 5mm or more and 12mm at the maximum, and has a tensile strength of 650MPa or more and a ratio of fatigue limit to yield strength (fatigue limit/yield strength) of 0.25 or more, so that high strength and excellent durability can be ensured.
Hereinafter, a method for manufacturing a high-strength steel material having excellent durability according to another aspect of the present invention will be described in detail.
The high-strength steel of the present invention can be produced by subjecting a steel slab satisfying the alloy composition proposed in the present invention to a series of processes of [ heating-hot rolling-cooling ].
Hereinafter, each process condition will be described in detail.
Heating steel billet
In the present invention, the process of heating the slab to homogenize the slab before hot rolling is preferably performed, and in this case, the heating process is preferably performed at 1200 to 1350 ℃.
When the heating temperature is less than 1200 ℃, precipitates cannot be sufficiently re-dissolved, so that the formation of precipitates is reduced in the process after hot rolling, and there is a problem that coarse TiN remains. On the other hand, when the heating temperature exceeds 1350 ℃, strength is lowered due to abnormal grain growth of austenite grains, which is not preferable.
Hot rolling
The reheated slab is preferably hot-rolled to produce a hot-rolled steel sheet, and at this time, it is preferably hot-finish rolled at a temperature in the range of 800 to 1150 ℃ under conditions satisfying the following [ relational expression 1 ].
When the hot rolling is performed at a temperature higher than 1150 deg.c, the temperature of the hot rolled steel sheet becomes high, resulting in coarsening of crystal grain size and deterioration of surface quality of the hot rolled steel sheet. On the other hand, when hot rolling is performed at a temperature lower than 800 ℃, the anisotropy becomes severe and the formability also becomes poor due to the development of extended crystal grains due to excessively delayed recrystallization, and the inhomogeneous fine structure becomes more developed due to rolling at a temperature lower than the austenite temperature region.
In particular, in the hot rolling process of the present invention, when the rolling is terminated at a temperature (temperature exceeding Tn) higher than the temperature range set forth in the following relational expression 1, the microstructure of the steel is coarse and non-uniform, and coarse MA phase and martensite phase are formed due to the delay of transformation, resulting in excessive microcracks being formed at the time of shear forming and punch forming, and thus the durability is deteriorated. On the other hand, when the rolling is terminated at a temperature lower than the temperature range (temperature lower than Tn-50) suggested in the following relational expression 1, in a thick steel material having a thickness of 5mm or more, ferrite transformation is promoted at a t/4 position in the thickness direction directly below the surface layer having a relatively low temperature to increase the phase fraction of fine ferrite, but the steel material has an extended grain shape to cause rapid crack propagation, and a non-uniform fine structure may remain in the thickness center portion, which is disadvantageous in securing durability.
[ relational expression 1]
FDT (hot rolling termination temperature (DEG C)) with Tn-50 (more than or equal to) and less than or equal to Tn
Tn =730+92 x [ C ] +70 x [ Mn ] +45 x [ Cr ] +650 x [ Nb ] +410 x [ Ti ] -80 x [ Si ] -1.4 x (t-5) (where the elements represent their weight percentages and t represents the thickness (mm) of the final hot-rolled steel sheet).
Cooling and winding
The hot rolled steel sheet manufactured by hot rolling as described above is cooled to a temperature ranging from 400 to 500 c, and then a coiling process may be performed at the temperature.
The cooling is performed by primary cooling and secondary cooling, and the primary cooling is preferably performed so as to satisfy [ relational expression 2]]Cooling Rate (CR) 1 ) The secondary cooling is preferably performed so as to satisfy the following [ relational expression 3]]Cooling Rate (CR) 2 ) The process is carried out.
Specifically, it is preferable to terminate the primary cooling in a temperature zone where a ferrite transformation occurs during cooling, but the temperature at which the ferrite transformation occurs may vary according to the alloy composition proposed in the present invention. More specifically, the primary cooling is preferably performed to a temperature at which the transformation of hard phases such as bainite phase, MA phase, martensite phase, etc. does not occur. Further preferably, the primary cooling may be performed until the temperature of the hot rolled steel sheet obtained by hot rolling reaches 600 ℃.
When the primary cooling is performed in the temperature range, as in the present invention, when the thickness of the rolled sheet is 5mm or more, the cooling rate in the central portion of the thickness of the rolled sheet is slower than the cooling rate in the region from immediately below the surface layer to t/4, and therefore, a coarse ferrite phase may be formed in the central portion of the thickness, and a nonuniform fine structure may be obtained.
Therefore, in the present invention, it is preferable to use a specific Cooling Rate (CR) represented by the following relational expression 2 1 ) The cooling is performed at a high cooling rate so that excessive formation of a ferrite phase or coarsening of the ferrite phase does not occur in the primary cooling process.
[ relational expression 2]
CR 1 ≥196-300×[C]+4.5×[Si]-71.8×[Mn]-59.6×[Cr]+187×[Ti]+852×[Nb](wherein each element represents a weight percentage thereof).
The secondary cooling is performed immediately after the primary cooling is terminated according to the above conditions, and is preferably terminated at a take-up temperature (CT (° c)).
In the secondary cooling in the temperature range, in order to convert the non-transformed phase into the bainite phase throughout the entire thickness of the steel material so that 90% (area fraction) of the matrix structure is formed of the ferrite phase and the bainite phase, a specific Cooling Rate (CR) represented by the following relational expression 3 is preferably used 2 ) And cooling is carried out. At this time, when the cooling rate ratio CR is set Minimum size When slow, carbides are formed and grown coarsely without forming a bainite phase, and the carbides exist mainly in grain boundaries of a ferrite phase, and when the cooling rate is slower, a pearlite phase is formed, so that cracks are easily formed at the time of shear forming or punch forming, and the cracks propagate along the grain boundaries even under a small external force. On the other hand, when the cooling rate exceeds CR Maximum of In the case of the phase (phase), the MA phase or the martensite phase is excessively formed to increase the difference in hardness between the phases, and thus the durability is deteriorated.
Therefore, when the secondary cooling is performed in the temperature range, the cooling needs to be performed at a cooling rate satisfying the following relational expression 3.
[ relational expression 3]
CR Minimum size of ≤CR 2 ≤CR Maximum of
(CR Maximum of ＝76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[Nb]，CR Minimum size ＝27.4-45.3×[C]+5.28×[Si]-11×[Mn]-7.33×[Cr]+42.3×[Ti]+82×[Nb]Each element represents a weight percentage thereof).
In addition, when the winding is performed after the completion of the above-mentioned cooling process, when the winding temperature exceeds 500 ℃, a pearlite phase is formed and thus the strength of the steel is insufficient, and on the other hand, when the winding temperature is less than 400 ℃, an excessive martensite phase is formed and thus the shear formability, the punch formability and the durability are deteriorated.
In the process of manufacturing the steel material desired in the present invention, the process conditions are controlled to satisfy the above relational expressions 1 to 3, so that the area fraction of the crystal grains having an aspect ratio of 0.3 or less formed in the thickness center portion of the steel material is ensured to be less than 50%, and the area per unit area (1 mm) can be ensured 2 ) The length of the grain boundary observed inside is 700mm or more.
In the process of manufacturing a thick steel material having a thickness of 5mm or more, it is difficult to uniformly secure a fine structure in the thickness center portion when performing a conventional hot rolling. In particular, in order to obtain a retardation effect of recrystallization in the thickness center portion, when hot rolling is performed at an excessively low temperature, a deformed structure develops strongly in a region from immediately below the surface layer to t/4 in the thickness direction of the rolled sheet, and phase unevenness with the thickness center portion increases, so that fine cracks are likely to occur in uneven portions during shear deformation or punching deformation, and the durability of the member is also deteriorated. Therefore, as shown in the relational expression 1, it is necessary to complete hot rolling between the Tn (. Degree. C.) temperature, which is the starting temperature of the delay of recrystallization, and Tn-50 (. Degree. C.).
When the hot rolling is terminated at a temperature higher than the temperature set forth in the above relational expression 1, the area fraction of crystal grains having an aspect ratio of 0.3 or less is greatly reduced due to the formation of coarse ferrite phases and polygonal ferrite phases, while the central portion strength may be reduced due to the significant reduction in the size of grain boundaries, and there is a problem that cracks are likely to grow when they are formed. Further, when the hot rolling is terminated at a temperature lower than the temperature set forth in the above-mentioned relational expression 1, the area fraction of the crystal grains having an aspect ratio of 0.3 or less is greatly increased due to the increase of the excessively elongated crystal grains, and cracks formed at the time of shear forming are easily propagated by an external force due to the formation of coarse carbide or martensite phase at the grain boundary, so that the durability is deteriorated.
In addition, the relational expressions 2 and 3 correspond to cooling conditions for optimizing the fine structure so that the strength and durability of the steel can be improved through a transformation process in cooling. That is, since the type and fraction of the structure phase change and the aspect ratio of crystal grains and the size of grain boundaries change depending on the cooling conditions, it is preferable to perform cooling under the conditions satisfying relational expressions 2 and 3.
Air cooling
The coil obtained by performing the cooling and winding processes as described above may be air-cooled to a temperature range of normal temperature to 200 ℃. In this case, the air cooling process of the coil means cooling in the atmosphere at a cooling rate of 0.001 to 10 ℃/hr. At this time, when the cooling rate exceeds 10 ℃/hr, a part of the non-transformed phase in the steel is easily transformed into the MA phase, and thus the shear formability, punch formability and durability of the steel are deteriorated. On the other hand, in order to control the cooling rate to less than 0.001 ℃/hr, additional heating equipment, heat-compensating equipment, and the like are required, which is economically disadvantageous.
In addition, the steel material air-cooled as described above is pickled and oiled, and then may be heated to a temperature range of 450 to 740 ℃ to perform a hot-dip galvanizing process.
The hot dip galvanizing process may use a zinc-based plating bath, and the alloy composition in the zinc-based plating bath is not particularly limited, but may be, as one example, a zinc alloy containing magnesium (Mg): 0.01-30 wt%, aluminum (Al): 0.01-50 wt% and the balance of Zn and unavoidable impurities.
The present invention will be described in more detail with reference to examples. It should be noted, however, that the following examples are only for illustrating the present invention and are not intended to limit the scope of the present invention. This is because the scope of the right of the present invention is determined by the contents recited in the claims and reasonably derived therefrom.
Detailed Description
(examples)
A steel slab having an alloy composition of table 1 below was prepared. At this time, the content of the alloy composition is wt%, and the balance includes Fe and inevitable impurities. The prepared billets were manufactured into steel materials according to the manufacturing conditions of table 2 below. At this time, when cooling is performed after hot rolling, primary cooling is completed at 600 ℃, and secondary cooling is completed at a coiling temperature.
In table 2 below, FDT represents the temperature at the time of finish hot rolling (hot rolling finish temperature), CT represents the coiling temperature, and a constant cooling rate of 1 ℃/hr is applied when air cooling is performed after coiling.
[ Table 1]
(the alloy compositions of comparative steels 1 to 11 in the above Table 1 satisfy the scope of the present invention, but the manufacturing conditions in the following Table 2 are not within the scope of the present invention and are therefore labeled as comparative steels.)
[ Table 2]
For each of the steel sheets manufactured as described above, mechanical properties such as Tensile Strength (TS), yield Strength (YS), and elongation (T-El) and durability were evaluated, and a fine structure was observed, and the results thereof are shown in table 3 below.
Specifically, the yield strength and elongation represent 0.2% offset (off-set) yield strength and elongation at break, respectively, and at the same time, the tensile strength was measured by a test piece of JIS5 standard, which is a test piece taken in a direction perpendicular to the rolling direction.
The durability was evaluated by subjecting a test piece having a punched portion to a high frequency fatigue test (bending fatigue test), and the results are shown. In this case, a test piece for fatigue test was produced by punching a hole having a diameter of 10mm at a clearance (clearance) of 12% in the center of a bending fatigue test piece having a gauge length of 40mm and a width of 20mm by a punch forming method, and testing the bending fatigue test piece under a stress ratio of-1 and a frequency of 15 Hz. Fatigue strength (S) Fatigue ) Is compared with yield strength and in strength ratio (S) Fatigue YS), from which changes in the cross-sectional quality and durability of the punched part could be confirmed.
The microstructure of each steel material is shown as a result observed at the center (t/2) in the thickness direction. Per unit area (1 mm) corresponding to the area of grain boundary 2 ) The grain boundary length of (a) and the Aspect Ratio (AR) of the crystal grains, and the crystal grains having a large angle grain boundary of 15 ° or more were measured by Electron Back Scattering Diffraction (EBSD). The area fraction of the MA phase is shown by the results of analysis at 1000 magnifications using an optical microscope and an Image analyzer (Image analyzer) after etching by the Lepera etching method. In addition, the phase fractions of martensite (M), ferrite (F), bainite (B), and pearlite (P) are measured from the results of analysis at 3000 magnifications and 5000 magnifications using a Scanning Electron Microscope (SEM).
In the following table 3, F denotes Polygonal Ferrite (Polygonal Ferrite) having an equiaxed grain shape, and B denotes the sum of all bainite phases and fractions of Ferrite phases such as acicular Ferrite, bainitic Ferrite, etc., observed in a low temperature region.
In addition, AR 0.3 in table 3 below represents the proportion (area fraction) of crystal grains having an aspect ratio of 0.3 or less, which is a result observed and obtained at 1000 magnifications.
[ Table 3]
As shown in tables 1 to 3, the matrix structures of inventive steels 1 to 7 satisfying all the alloy compositions and manufacturing conditions proposed in the present invention were formed of ferrite and bainite composite structures. Further, the fraction of crystal grains having an aspect ratio of 0.3 or less in the central portion in the thickness direction of the steel material is less than 50% (refer to fig. 2), and all the grain boundary lengths are also formed to 700mm or more, thus ensuring desired high strength and excellent durability.
On the other hand, comparative steels 1 to 11 are the ones that satisfy the alloy compositions proposed in the present invention, but the production conditions are not within the range of the present invention, and thus the desired physical properties cannot be secured.
On the other hand, comparative steels 2 and 3 are thick steel materials having a thickness of 10mm and 7mm, respectively, and since the MA phase is developed in the microstructure of the central portion of comparative steel 2 and the length of the grain boundary is formed to be less than 700mm, when exposed to a fatigue environment, micro cracks formed in the cross section are likely to grow, and poor fatigue characteristics are exhibited. In comparative steel 3, since hot rolling in a low temperature range causes excessively elongated crystal grains to form in the thickness center portion, it is judged that fatigue fracture occurs along the brittle grain boundary. That is, this is because fine cracks develop along the ferrite grain boundary extending in the thickness center portion during punching.
In the case where comparative steel 10 does not satisfy relational expression 3, that is, the cooling rate at the time of secondary cooling is not within the range of the present invention, it is judged that when exposed to a fatigue environment, fracture easily occurs in a region where the difference in hardness from the surrounding phase (phase) is large, because the cooling rate at the time of secondary cooling is too high and the martensite phase is excessively formed in the thickness center portion.
In addition, comparative steels 12 to 17 were cases where the alloy compositions were out of the range of the present invention, and all of them satisfied relational expressions 1 to 3 at the time of production, and had the same thickness (8 mm), but had poor fatigue characteristics.
Specifically, the comparative steel 12 is a case where the C content is insufficient, an excessive ferrite phase is formed in the thickness center portion, and a bainite phase is not sufficiently formed. This coarsens the fine structure and lowers the fatigue strength.
In comparative steel 13, in which the content of C was too large, during transformation, an excessive pearlite phase and martensite phase were formed due to the high C concentration in the untransformed phase, and thus low fatigue strength was exhibited with respect to yield strength.
In comparative steel 14, the Si content was too high, the MA phase was formed together with the bainite phase, and a large amount of elongated microstructure was observed. Therefore, the fatigue characteristics were poor, and it was judged that this was caused by the formation of a large number of cracks around the MA phase, which is a relatively hard phase.
Comparative steel 15 was insufficient in Mn content, and although it was produced so as to satisfy relational expressions 1 to 3 in order to obtain a recrystallization retardation effect and a uniform microstructure, a ferrite phase was excessively formed in the thickness center portion, and thus low strength and fatigue strength were exhibited.
In comparative steel 16, the Mn content was too high, and the martensite phase was excessively developed along the Mn segregation band developed in the thickness center portion, and therefore, the cross-sectional quality and the fatigue property were inferior.
In addition, in the case of comparative steel 17 containing an excessive amount of Cr, a large amount of martensite phase formed locally was observed in the thickness center portion similarly to the case of comparative steel 16, and thus the fatigue characteristics were poor.
In addition, the fine structure of inventive steel 4 was observed, and then the results of measuring the aspect ratio of the crystal grains are shown in table 4 below.
[ Table 4]
In table 4, the minimum value and the maximum value represent the minimum value and the maximum value of the aspect ratio (the short side length of the crystal grain/the long side length of the crystal grain), respectively, and the Total Fraction (Total Fraction) represents the area Fraction of the crystal grain corresponding to a range exceeding the minimum value to the maximum value or less.
As shown in table 4, in the case of invention steel 4, it was confirmed that the fraction of crystal grains having an aspect ratio (short side/long side ratio) of 0.3 or less was less than 50% (total fraction was less than 0.5).
Claims (10)
1. A high-strength steel material having excellent durability, comprising, in weight%: carbon (C): 0.05-0.15%, silicon (Si): 0.01-1.0%, manganese (Mn): 1.0-2.3%, aluminum (Al): 0.01-0.1%, chromium (Cr): 0.005-1.0%, phosphorus (P): 0.001-0.05%, sulfur (S): 0.001-0.01%, nitrogen (N): 0.001-0.01%, niobium (Nb): 0.005-0.07%, titanium (Ti): 0.005-0.11%, and the balance of Fe and other unavoidable impurities,
the fine structure has a sum of area fractions of a ferrite phase and a bainite phase of 90% or more, and a fraction of crystal grains having an aspect ratio of 0.3 or less in a central portion of less than 50%, the central portion being a position from t/4 to a position t/2 in a thickness direction, the aspect ratio being a ratio of short sides to long sides, the central portion having an area of 1mm 2 The length of the grain boundary observed therein is 700mm or more.
2. The high-strength steel material excellent in durability according to claim 1, wherein an area fraction of a mixed structure of martensite and austenite of the steel material is less than 3%.
3. The high-strength steel material excellent in durability according to claim 1, wherein the steel material has a sum of an area fraction of a martensite phase and a mixed structure of martensite and austenite of 1 to 10%.
4. The high-strength steel material having excellent durability according to claim 1, wherein the steel material has a tensile strength of 650MPa or more and a fatigue limit/yield strength ratio of a fatigue limit to a yield strength of 0.25 or more.
5. A method of producing a high-strength steel material excellent in durability according to claim 1, characterized by comprising the steps of:
heating a steel slab at a temperature in the range of 1200-1350 ℃, said steel slab comprising in weight-%: carbon (C): 0.05-0.15%, silicon (Si): 0.01-1.0%, manganese (Mn): 1.0-2.3%, aluminum (Al): 0.01-0.1%, chromium (Cr): 0.005-1.0%, phosphorus (P): 0.001-0.05%, sulfur (S): 0.001-0.01%, nitrogen (N): 0.001-0.01%, niobium (Nb): 0.005-0.07%, titanium (Ti): 0.005-0.11%, and the balance of Fe and other unavoidable impurities;
hot rolling the heated slab to manufacture a hot rolled steel sheet;
cooling the hot rolled steel plate to the temperature range of 400-500 ℃ and then rolling (CT); and
after the rolling, the steel wire is air-cooled to the temperature range from normal temperature to 200 ℃,
wherein the hot rolling is a finish hot rolling performed at a temperature FDT satisfying [ relational formula 1] below in units of,
the cooling is performed with primary cooling to satisfy [ relational expression 2] and secondary cooling]Cooling rate CR of 1 Is carried out so as to satisfy the following [ relational expression 3]]Cooling rate CR of 2 The process is carried out by the following steps,
[ relational expression 1]
The hot rolling termination temperature FDT is more than or equal to Tn-50 and less than or equal to Tn
Tn＝730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]-80×[Si]-1.4×(t-5)
Wherein each element represents the weight percentage thereof, t represents the thickness mm of the final hot rolled steel sheet,
[ relational expression 2]
CR 1 ≥196-300×[C]+4.5×[Si]-71.8×[Mn]-59.6×[Cr]+187×[Ti]+852×[Nb]
Wherein, each element represents the weight percentage content,
[ relational expression 3]
CR Minimum size ≤CR 2 ≤CR Maximum of
CR Maximum of ＝76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[Nb]，CR Minimum size ＝27.4-45.3×[C]+5.28×[Si]-11×[Mn]-7.33×[Cr]+42.3×[Ti]+82×[Nb]Each element represents the weight percentage thereof.
6. The method for producing a high-strength steel material excellent in durability according to claim 5, wherein the primary cooling is terminated at 600 ℃.
7. The method for producing a high-strength steel material excellent in durability according to claim 5, wherein the secondary cooling is terminated at a winding temperature CT in units of ℃.
8. The method for producing a high-strength steel material excellent in durability according to claim 5, further comprising a step of pickling and oiling the steel sheet after the cooling.
9. The method of manufacturing a high-strength steel material excellent in durability according to claim 8, further comprising a step of hot-dip galvanizing after heating the steel sheet to a temperature range of 450 to 740 ℃ after the pickling and oiling.
10. The method for producing a high-strength steel material excellent in durability according to claim 9, wherein a plating bath used in the hot-dip galvanizing includes: magnesium (Mg): 0.01-30 wt%, aluminum (Al): 0.01-50% and the balance of Zn and unavoidable impurities.
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| PCT/KR2019/016299 WO2020111702A1 (en) | 2018-11-26 | 2019-11-26 | High-strength steel with excellent durability and method for manufacturing same |
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| KR102409896B1 (en) * | 2020-10-23 | 2022-06-20 | 주식회사 포스코 | High strength steel plate having excellent workability and method for manufacturing the same |
| KR102403648B1 (en) * | 2020-11-17 | 2022-05-30 | 주식회사 포스코 | High strength hot-rolled steel sheet and hot-rolled plated steel sheet, and manufacturing method for thereof |
| CN113084453B (en) * | 2021-03-18 | 2022-04-29 | 湖南三一路面机械有限公司 | Wear-resistant steel wheel, steel wheel manufacturing method, steel wheel welding method and road roller |
| KR20230072050A (en) | 2021-11-17 | 2023-05-24 | 주식회사 포스코 | High strength steel plate having excellent impact toughness after cold forming and high yield ratio and method for manufacturing the same |
| KR20240098898A (en) * | 2022-12-21 | 2024-06-28 | 주식회사 포스코 | Hot rolled steel sheet and method for the same |
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