WO2010030021A1 - 高強度鋼板およびその製造方法 - Google Patents
高強度鋼板およびその製造方法 Download PDFInfo
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- WO2010030021A1 WO2010030021A1 PCT/JP2009/065981 JP2009065981W WO2010030021A1 WO 2010030021 A1 WO2010030021 A1 WO 2010030021A1 JP 2009065981 W JP2009065981 W JP 2009065981W WO 2010030021 A1 WO2010030021 A1 WO 2010030021A1
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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
- 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
- C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
- C21D1/20—Isothermal quenching, e.g. bainitic hardening
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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
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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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- 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
- C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
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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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- 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
- C22C38/002—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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- 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
- C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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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
- C22C38/02—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
- 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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/001—Austenite
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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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/002—Bainite
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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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
Definitions
- the present invention relates to a high-strength steel sheet having a tensile strength (TS) of 980 MPa or more excellent in workability, particularly ductility and stretch flangeability, used in industrial fields such as automobiles and electrical equipment, and a method for producing the same.
- TS tensile strength
- the workability of the steel plate is strongly influenced by the workability of the hard phase. This is because when the ratio of hard phase is small and soft polygonal ferrite is large, the deformability of polygonal ferrite dominates the workability of the steel sheet, and even when the hard phase has insufficient workability. Workability such as ductility was ensured, but when the ratio of the hard phase is large, the deformability of the hard phase itself directly affects the formability of the steel sheet, not the deformation of polygonal ferrite, and the hard phase itself This is because if the workability is insufficient, the workability of the steel sheet is significantly deteriorated.
- steel plates having a hard phase other than martensite there are steel plates in which the main phase is polygonal ferrite, the hard phase is bainite or pearlite, and carbides are generated in these hard phases bainite or pearlite.
- This steel sheet is a steel sheet that not only improves the workability with polygonal ferrite alone, but also improves the workability of the hard phase itself by generating carbides in the hard phase, and in particular, improves the stretch flangeability. .
- the main phase is polygonal ferrite, it is difficult to achieve both high strength and workability of 980 MPa or higher in tensile strength (TS).
- Patent Document 1 proposes a high-tensile steel plate that is excellent in bending workability and impact properties by defining alloy components and making the steel structure fine and uniform bainite having retained austenite.
- Patent Document 2 proposes a composite structure steel plate having excellent bake hardenability by defining predetermined alloy components, making the steel structure bainite having retained austenite, and defining the amount of retained austenite in bainite. ing.
- Patent Document 3 a predetermined alloy component is defined, the steel structure is 90% or more in area ratio of bainite having retained austenite, the amount of retained austenite in bainite is 1% or more and 15% or less, and the hardness of bainite.
- HV a composite structure steel plate excellent in impact resistance
- the above-described steel sheet has the following problems.
- it is difficult to ensure a stable amount of retained austenite that exhibits the TRIP effect in a high strain region when strain is applied to the steel sheet, and bendability is obtained.
- the ductility until plastic instability occurs is low, and the stretchability is inferior.
- Patent Document 2 has bake hardenability, even when trying to increase the tensile strength (TS) to 980 MPa or higher, or even 1050 MPa or higher, bainite or martensite mainly composed of ferrite is suppressed as much as possible. Therefore, it is difficult to ensure workability such as ductility and stretch flangeability at the time of securing strength or increasing strength.
- TS tensile strength
- the main purpose of the steel sheet described in Patent Document 3 is to improve impact resistance, and since it has a main phase of bainite having a hardness of HV 250 or less, specifically, it has a structure containing more than 90%. It is difficult to set the tensile strength (TS) to 980 MPa or more.
- the present invention advantageously solves the above-mentioned problems, and provides a high-strength steel sheet having a tensile strength (TS) of 980 MPa or more, which is excellent in workability, particularly ductility and stretch flangeability, together with its advantageous production method.
- the high-strength steel sheet of the present invention includes a steel sheet obtained by subjecting the surface of the steel sheet to hot dip galvanization or galvannealing.
- excellent workability means that TS ⁇ T.
- the value of EL satisfies 20000 MPa ⁇ % or more and the value of TS ⁇ ⁇ satisfies 25000 MPa ⁇ %.
- TS is tensile strength (MPa)
- T.I. EL is the total elongation (%)
- ⁇ is the critical hole expansion rate (%).
- the inventors have made extensive studies on the component composition and microstructure of the steel sheet in order to solve the above problems.
- the lower bainite structure and / or martensite structure is utilized to increase the strength, and the C content in the steel sheet is increased to 0.17% or more and the C content is increased, and then the upper bainite transformation is utilized.
- stable retained austenite advantageous for obtaining the TRIP effect can be secured, and by making a part of the martensite tempered martensite, workability, in particular, balance between strength and ductility, and strength and stretch flange It was found that a high-strength steel sheet having a tensile strength of 980 MPa or more that is excellent in both the balance of properties can be obtained.
- the present invention is based on the above findings, and the gist of the present invention is as follows. 1. C: 0.17% to 0.73% by mass%, Si: 3.0% or less, Mn: 0.5% to 3.0%, P: 0.1% or less, S: 0.07% or less, Al: 3.0% or less and N: 0.010% or less, and Si + Al satisfies 0.7% or more, the balance is the composition of Fe and inevitable impurities, As the steel sheet structure, the area ratio of the total amount of lower bainite and all martensite is 10% or more and 90% or less, the amount of retained austenite is 5% or more and 50% or less, and the steel structure of bainitic ferrite in the upper bainite.
- the area ratio with respect to the whole is 5% or more, of the total amount of the lower bainite and all martensite, 75% or less of the as-quenched martensite, and the area ratio of the polygonal ferrite with respect to the entire steel sheet structure is 10% or less (0% A high-strength steel sheet characterized in that the average C content in the retained austenite is 0.70% or more and the tensile strength is 980 MPa or more.
- the steel sheet is further in mass%, Cr: 0.05% or more and 5.0% or less, 2.
- the steel sheet is further in mass%, 1 or 2 above, which contains one or two elements selected from Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%. High strength steel sheet as described.
- the steel sheet is further in mass%
- B The high-strength steel sheet according to any one of 1 to 3 above, which contains 0.0003% or more and 0.0050% or less.
- the steel sheet is further in mass%, 1 to 4 above, which contains one or two elements selected from Ni: 0.05% to 2.0% and Cu: 0.05% to 2.0%
- the high-strength steel sheet according to any one of the items.
- the steel sheet is further in mass%, 1 to 5 above, which contains one or two elements selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%.
- the high-strength steel sheet according to any one of the items.
- a high-strength steel sheet comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel sheet according to any one of 1 to 6 above.
- the steel slab having the composition described in any one of 1 to 6 above is hot-rolled to form a cold-rolled steel sheet by cold rolling, and then the cold-rolled steel sheet is 600 seconds or more in the austenite single-phase region.
- cooling stop temperature determined in the first temperature range of 350 ° C. or higher and 490 ° C. or lower: when cooling to T ° C., the average cooling rate is controlled to 5 ° C./s or higher until at least 550 °
- the high-strength steel sheet is cooled and then held in the first temperature range for 15 seconds to 1000 seconds and then held in the second temperature range of 200 ° C. to 350 ° C. for 15 seconds to 1000 seconds. Manufacturing method.
- TS tensile strength
- the utility value is extremely large, and is particularly useful for reducing the weight of automobile bodies.
- the area ratio is the area ratio relative to the entire steel sheet structure.
- Area ratio of total amount of lower bainite and all martensite 10% or more and 90% or less
- Lower bainite and martensite are structures necessary for increasing the strength of a steel sheet. If the area ratio of the total amount of lower bainite and all martensite is less than 10%, the tensile strength (TS) of the steel sheet does not satisfy 980 MPa. On the other hand, if the area ratio of the total amount of the lower bainite and all martensite exceeds 90%, the upper bainite decreases, and as a result, stable retained austenite enriched in C cannot be secured, so workability such as ductility is reduced. Decreasing becomes a problem. Therefore, the area ratio of the total amount of lower bainite and all martensite is set to 10% or more and 90% or less. Preferably, it is in the range of 20% to 80%. More preferably, it is in the range of 30% to 70%.
- ratio of as-quenched martensite 75% or less
- the ratio of as-quenched martensite is the sum of lower bainite and all martensite present in the steel sheet. If it exceeds 75% of the amount, the tensile strength becomes 980 MPa or more, but the stretch flangeability is inferior.
- the as-quenched martensite is extremely hard and the deformability of the as-quenched martensite itself is extremely low, so that the workability of the steel sheet, particularly the stretch flangeability, is significantly deteriorated.
- the ratio of martensite as quenched in martensite is 75% or less with respect to the total amount of lower bainite and all martensite present in the steel sheet. Preferably it is 50% or less.
- the as-quenched martensite is a structure in which carbides are not recognized in the martensite and can be observed by SEM.
- Residual austenite amount 5% or more and 50% or less Residual austenite undergoes martensitic transformation by the TRIP effect during processing, and improves ductility by increasing strain dispersibility.
- Residual austenite amount 5% or more and 50% or less Residual austenite undergoes martensitic transformation by the TRIP effect during processing, and improves ductility by increasing strain dispersibility.
- utilizing the upper bainite transformation in particular, retained austenite with an increased amount of C concentration is formed in the upper bainite.
- retained austenite that can exhibit the TRIP effect even in a high strain region during processing can be obtained.
- good workability can be obtained even in a high strength region where the tensile strength (TS) is 980 MPa or more, specifically, TS ⁇ T.
- the value of El can be set to 20000 MPa or more, and a steel sheet having an excellent balance between strength and ductility can be obtained.
- the retained austenite in the upper bainite is formed between the laths of the bainitic ferrite in the upper bainite and is finely distributed. Is necessary and accurate quantification is difficult.
- the amount of retained austenite formed between the laths of the bainitic ferrite is a certain amount commensurate with the amount of bainitic ferrite formed.
- the area ratio of bainitic ferrite in the upper bainite is 5% or more
- X-ray diffraction is a method for measuring the amount of retained austenite that has been conventionally performed. If the amount of retained austenite obtained from the strength measurement, specifically the X-ray diffraction intensity ratio of ferrite and austenite is 5% or more, a sufficient TRIP effect can be obtained, and the tensile strength (TS) is 980 MPa or more. TS ⁇ T. It was found that El can achieve 20000 MPa ⁇ % or more.
- the amount of retained austenite obtained by a conventional method for measuring the amount of retained austenite is equivalent to the area ratio of retained austenite to the entire steel sheet structure.
- the amount of retained austenite is in the range of 5% to 50%.
- it is in the range of more than 5%, more preferably 10% or more and 45% or less. More preferably, it is the range of 15% or more and 40% or less.
- Average C content in retained austenite 0.70% or more
- TS tensile strength
- the amount of C in the austenite is important.
- C is concentrated in the retained austenite formed between the laths of bainitic ferrite in the upper bainite.
- the conventional austenite in the retained austenite If the average C content in the retained austenite obtained from the shift amount of the diffraction peak in X-ray diffraction (XRD), which is a method for measuring the average C content (average of the C content in the retained austenite) is 0.70% or more It was found that excellent processability can be obtained. When the average C content in the retained austenite is less than 0.70%, martensitic transformation occurs in the low strain region during processing, and the TRIP effect in the high strain region that improves workability cannot be obtained.
- XRD X-ray diffraction
- the average amount of C in the retained austenite is 0.70% or more. Preferably it is 0.90% or more.
- the average C content in the retained austenite is preferably 2.00% or less. More preferably, it is 1.50% or less.
- the area ratio of bainitic ferrite in the upper bainite 5% or more
- the formation of bainitic ferrite by the upper bainite transformation concentrates C in the untransformed austenite and exhibits the TRIP effect in the high strain region during processing. It is necessary to obtain retained austenite that enhances strain resolution.
- the transformation from austenite to bainite occurs over a wide temperature range of approximately 150 to 550 ° C., and various types of bainite are produced within this temperature range. In the prior art, such various bainite was often simply defined as bainite, but in order to obtain the target workability in the present invention, it is necessary to clearly define the bainite structure.
- the bainite and lower bainite are defined as follows.
- the upper bainite is composed of lath-like bainitic ferrite and residual austenite and / or carbide existing between bainitic ferrite, and there is no fine carbide regularly arranged in lath-like bainitic ferrite. It is a feature.
- the lower bainite is composed of the lath-shaped bainitic ferrite and the residual austenite and / or carbide existing between the bainitic ferrites in common with the upper bainite. It is characterized by the presence of fine carbides regularly arranged in the bainitic ferrite. That is, the upper bainite and the lower bainite are distinguished by the presence or absence of regularly arranged fine carbides in bainitic ferrite.
- Such a difference in the state of carbide formation in bainitic ferrite has a great influence on the concentration of C in the retained austenite. That is, when the area ratio of the bainitic ferrite of the upper bainite is less than 5%, even when the bainite transformation is advanced, the amount of C generated as carbides in the bainitic ferrite increases, resulting in a The amount of C enriched in the residual austenite present in the steel decreases, and the amount of residual austenite that exhibits the TRIP effect in the high strain region during processing decreases. Therefore, the area ratio of bainitic ferrite in the upper bainite needs to be 5% or more in terms of the area ratio with respect to the entire steel sheet structure. On the other hand, if the area ratio of the bainitic ferrite of the upper bainite to the entire steel sheet structure exceeds 85%, it may be difficult to ensure the strength.
- Polygonal ferrite area ratio 10% or less (including 0%)
- TS tensile strength
- the area ratio of polygonal ferrite is 10% or less, even if polygonal ferrite is present, a small amount of polygonal ferrite is isolated and dispersed in the hard phase, and strain concentration can be suppressed. Degradation of workability can be avoided. Therefore, the area ratio of polygonal ferrite is 10% or less. Preferably it is 5% or less, More preferably, it is 3% or less, and 0% may be sufficient.
- the hardness of the hardest structure in the steel sheet structure is HV ⁇ 800. That is, in the steel sheet of the present invention, when there is no as-quenched martensite, either the tempered martensite or the lower bainite or the upper bainite is the hardest phase, but these structures all have HV ⁇ 800. It is a phase. In addition, when there is as-quenched martensite, the as-quenched martensite becomes the hardest structure, in the steel sheet of the present invention, even if it is as-quenched martensite, the hardness is HV ⁇ 800, There is no extremely hard martensite such that HV> 800, and good stretch flangeability can be secured.
- the steel sheet of the present invention may contain pearlite, Widmanstatten ferrite, or lower bainite as the remaining structure.
- the allowable content of the remaining tissue is preferably 20% or less in terms of area ratio. More preferably, it is 10% or less.
- C 0.17% or more and 0.73% or less
- C is an element indispensable for increasing the strength of a steel sheet and ensuring a stable retained austenite amount, and for ensuring the amount of martensite and allowing austenite to remain at room temperature. It is a necessary element. If the C content is less than 0.17%, it is difficult to ensure the strength and workability of the steel sheet. On the other hand, if the amount of C exceeds 0.73%, the welded part and the heat-affected zone are hardened and the weldability deteriorates. Accordingly, the C content is in the range of 0.17% to 0.73%. Preferably, it is 0.20% or more and 0.48% or less of range, More preferably, it is 0.25% or more.
- Si 3.0% or less (including 0%) Si is a useful element that contributes to improving the strength of steel by solid solution strengthening. However, if the amount of Si exceeds 3.0%, the workability and toughness deteriorate due to the increase in the amount of solid solution in polygonal ferrite and bainitic ferrite, and the surface properties due to the occurrence of red scale, etc. In the case of deterioration or hot dipping, it causes deterioration of plating adhesion and adhesion. Therefore, the Si content is 3.0% or less. Preferably it is 2.6% or less. More preferably, it is 2.2% or less. Si is an element useful for suppressing the formation of carbides and promoting the formation of retained austenite. Therefore, the Si content is preferably 0.5% or more, but the formation of carbides is only Al. In the case of suppressing by Si, Si does not need to be added, and the Si amount may be 0%.
- Mn 0.5% or more and 3.0% or less Mn is an element effective for strengthening steel. If the amount of Mn is less than 0.5%, carbide precipitates in a temperature range higher than the temperature at which bainite and martensite are generated during cooling after annealing, so ensure the amount of hard phase that contributes to strengthening of steel. I can't. On the other hand, when the amount of Mn exceeds 3.0%, castability is deteriorated. Accordingly, the amount of Mn is set in the range of 0.5% to 3.0%. Preferably, the range is 1.5% or more and 2.5% or less.
- P 0.1% or less
- P is an element useful for strengthening steel, but if the P content exceeds 0.1%, the impact resistance deteriorates due to embrittlement due to grain boundary segregation, and the steel is alloyed. In the case of applying hot dip galvanizing, the alloying speed is greatly delayed. Therefore, the P content is 0.1% or less. Preferably it is 0.05% or less.
- the amount of P is preferably reduced, but if it is less than 0.005%, it causes a significant increase in cost, so the lower limit is preferably about 0.005%.
- S 0.07% or less Since S generates MnS and becomes inclusions, which causes deterioration of impact resistance and cracks along the metal flow of the weld, it is preferable to reduce the amount of S as much as possible. However, excessively reducing the amount of S causes an increase in manufacturing cost, so the amount of S is set to 0.07% or less. Preferably it is 0.05% or less, More preferably, it is 0.01% or less. In addition, since it is accompanied by a big increase in manufacturing cost to make S less than 0.0005%, the lower limit is about 0.0005% from the point of manufacturing cost.
- Al 3.0% or less
- Al is a useful element to be added as a deoxidizer in the steel making process, as well as a useful element for strengthening steel.
- the Al content is 3.0% or less.
- Al is an element useful for suppressing the formation of carbides and promoting the formation of retained austenite.
- the Al content should be 0.001% or more.
- it is more preferably 0.005% or more.
- the amount of Al in the present invention is the amount of Al contained in the steel sheet after deoxidation.
- N 0.010% or less
- N is an element that greatly deteriorates the aging resistance of steel, and is preferably reduced as much as possible.
- the N content exceeds 0.010%, deterioration of aging resistance becomes remarkable, so the N content is set to 0.010% or less. Note that, if N is less than 0.001%, a large increase in manufacturing cost is caused, so that the lower limit is about 0.001% from the viewpoint of manufacturing cost.
- the basic component has been described above. However, in the present invention, it is not sufficient to satisfy the above component range, and it is necessary to satisfy the following equation. Si + Al ⁇ 0.7% As described above, both Si and Al are useful elements for suppressing the formation of carbides and promoting the formation of retained austenite. Although suppression of the formation of carbides is effective even if Si or Al is contained alone, it is necessary to satisfy 0.7% or more in total of the Si amount and the Al amount.
- the amount of Al in the above formula is the amount of Al contained in the steel sheet after deoxidation.
- the component described below other than the above-mentioned basic component can be contained appropriately.
- One or more selected from Cr: 0.05% to 5.0%, V: 0.005% to 1.0% and Mo: 0.005% to 0.5% , V and Mo are elements having an action of suppressing the formation of pearlite during cooling from the annealing temperature. The effect is obtained when Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more.
- Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more if it exceeds Cr: 5.0%, V: 1.0% and Mo: 0.5%, the amount of hard martensite becomes excessive, and the strength becomes higher than necessary. Accordingly, when Cr, V and Mo are contained, Cr: 0.05% to 5.0%, V: 0.005% to 1.0% and Mo: 0.005% to 0.5% % Or less.
- Ti and Nb are useful for the precipitation strengthening of steel.
- Each content is 0.01% or more.
- the workability and the shape freezing property are lowered. Therefore, when Ti and Nb are contained, the range is Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%.
- B 0.0003% or more and 0.0050% or less B is an element useful for suppressing the formation and growth of ferrite from the austenite grain boundary. The effect is obtained when the content is 0.0003% or more. On the other hand, if the content exceeds 0.0050%, the workability decreases. Therefore, when it contains B, it is set as B: 0.0003% or more and 0.0050% or less of range.
- Ni and Cu are effective elements for strengthening steel. Moreover, when performing hot dip galvanization or alloying hot dip galvanization to a steel plate, the internal oxidation of a steel plate surface layer part is accelerated
- Ca and REM spheroidize the shape of the sulfide, and stretch flange Useful to improve the negative effects of sulfides on sex.
- the effect is obtained when each content is 0.001% or more.
- the respective contents exceed 0.005%, inclusions and the like increase, causing surface defects and internal defects. Therefore, when Ca and REM are contained, the range is Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%.
- components other than the above are Fe and inevitable impurities. However, as long as the effects of the present invention are not impaired, the inclusion of components other than those described above is not rejected.
- a steel slab adjusted to the above preferred component composition is manufactured, then hot-rolled, and then cold-rolled to obtain a cold-rolled steel sheet.
- these treatments are not particularly limited, and may be performed according to ordinary methods.
- the preferred production conditions are as follows. After heating the steel slab to a temperature range of 1000 ° C. or higher and 1300 ° C. or lower, hot rolling is finished in a temperature range of 870 ° C. or higher and 950 ° C. or lower, and the obtained hot rolled steel sheet is heated to a temperature of 350 ° C. or higher and 720 ° C. or lower. Take up in the area.
- the hot-rolled steel sheet is pickled and then cold-rolled at a rolling reduction in the range of 40% to 90% to obtain a cold-rolled steel sheet.
- the steel sheet is manufactured through normal steelmaking, casting, hot rolling, pickling and cold rolling processes.
- the steel plate is heated by thin slab casting or strip casting. You may manufacture by omitting a part or all of a hot rolling process.
- the obtained cold-rolled steel sheet is subjected to the heat treatment shown in FIG.
- Annealing is performed for 15 seconds to 600 seconds in an austenite single phase region.
- the steel sheet of the present invention is mainly composed of upper bainite, lower bainite and martensite which are transformed from untransformed austenite in a relatively low temperature range of 350 ° C. or more and 490 ° C. or less.
- annealing in the austenite single phase region is necessary.
- the annealing temperature is not particularly limited as long as it is in the austenite single phase region, but if the annealing temperature exceeds 1000 ° C., the growth of austenite grains is remarkable, which causes coarsening of the constituent phases caused by subsequent cooling, and deteriorates toughness and the like.
- the annealing temperature is of less than A 3 point (austenitic transformation point) is already generated by the polygonal ferrite in the annealing step, in order to suppress the growth of polygonal ferrite during cooling than 500 ° C. It becomes necessary to cool the temperature range very rapidly. Accordingly, the annealing temperature needs to be 3 points or more, and is preferably 1000 ° C. or less.
- the annealing time is in the range of 15 seconds to 600 seconds. Preferably, it is the range of 60 seconds or more and 500 seconds or less.
- [X%] is defined as mass% of the component element X of the steel sheet.
- the cold-rolled steel sheet after annealing is cooled to a cooling stop temperature: T ° C. determined in a first temperature range of 350 ° C. or more and 490 ° C. or less, but at least up to 550 ° C., the average cooling rate is controlled to 5 ° C./s or more. And cooled.
- the average cooling rate from the annealing temperature to the first temperature range is set to 5 ° C./s or more. Preferably, it is 10 ° C./s or more.
- the upper limit of the average cooling rate is not particularly limited as long as the cooling stop temperature does not vary, but in general equipment, when the average cooling rate exceeds 100 ° C./s, the structure in the longitudinal direction and the sheet width direction of the steel plate
- the dispersion is significantly increased, so that it is preferably 100 ° C./s or less.
- the steel sheet cooled to 550 ° C. is continuously cooled to a cooling stop temperature: T ° C.
- the speed at which the steel sheet is cooled in the temperature range of T ° C. or more and 550 ° C. or less is not particularly limited except that the holding time in the first holding temperature range is 15 seconds or more and 1000 seconds or less, but the steel sheet is excessively slow.
- the steel plate is preferably cooled at an average rate of 1 ° C./s or higher.
- Cooling stop temperature The steel sheet cooled to T ° C is held for a period of 15 seconds to 1000 seconds in a first temperature range of 350 ° C to 490 ° C.
- the upper limit of the first temperature range exceeds 490 ° C.
- carbide precipitates from untransformed austenite and a desired structure cannot be obtained.
- the lower limit of the first temperature range is less than 350 ° C.
- lower bainite is generated instead of upper bainite, and there is a problem that the amount of C concentration in austenite decreases. Therefore, the range of the first temperature range is 350 ° C. or more and 490 ° C. or less.
- the holding time in the first temperature range is less than 15 seconds, the amount of upper bainite transformation is reduced, and the amount of C enrichment in untransformed austenite is problematic.
- the retention time in the first temperature range exceeds 1000 seconds, stable retained austenite in which C is concentrated by precipitation of carbides from untransformed austenite which becomes retained austenite as the final structure of the steel sheet cannot be obtained.
- the desired processability cannot be obtained. Therefore, the holding time is 15 seconds or more and 1000 seconds or less. Preferably, it is the range of 30 seconds or more and 600 seconds or less.
- the steel sheet that has been held in the first temperature range is cooled to a second temperature range of 200 ° C. or higher and 350 ° C. or lower at an arbitrary rate, and held in the second temperature range for 15 seconds to 1000 seconds.
- the upper limit of the second temperature range exceeds 350 ° C.
- the lower bainite transformation does not proceed, resulting in a problem that the amount of martensite as quenched is increased.
- the lower limit of the second temperature range is less than 200 ° C., similarly, the lower bainite transformation does not proceed and the amount of martensite as quenched is increased. Therefore, the range of the second temperature range is 200 ° C. or more and 350 ° C. or less.
- the holding time is in the range of 15 seconds to 1000 seconds. Preferably, it is the range of 30 seconds or more and 600 seconds or less.
- the holding temperature does not need to be constant as long as it is within the predetermined temperature range described above, and even if it fluctuates within the predetermined temperature range, the gist of the present invention is not impaired.
- the cooling rate As long as the thermal history is satisfied, the steel sheet may be heat-treated with any equipment.
- the method for producing a high-strength steel sheet of the present invention can be further subjected to hot dip galvanizing treatment or galvannealing treatment obtained by adding alloying treatment to hot dip galvanizing treatment.
- the hot dip galvanizing process or the alloying hot dip galvanizing process may be performed during the cooling to the first temperature range or in the first temperature range.
- the holding time in the first temperature range is 15 seconds or more and 1000 seconds or less including the processing time in the first temperature range of the hot dip galvanizing process or the alloying galvanizing process.
- the hot dip galvanizing treatment or alloying hot dip galvanizing treatment is preferably performed in a continuous hot dip galvanizing line.
- the hot-dip galvanizing treatment or the alloying hot-dip galvanizing treatment is performed again. You can add that. Further, according to the production method of the present invention, the hot dip galvanizing treatment or the alloying hot dip galvanizing treatment can be performed after the holding in the second temperature range.
- the method of performing hot dip galvanizing treatment or alloying hot dip galvanizing treatment on a steel sheet is as follows.
- the steel sheet is infiltrated into the plating bath and the amount of adhesion is adjusted by gas wiping.
- the amount of dissolved Al in the plating bath ranges from 0.12% to 0.22% in the case of hot dip galvanizing, and ranges from 0.08% to 0.18% in the case of galvannealed alloying. It is preferable that In the case of hot dip galvanizing, the temperature of the plating bath may be in the range of 450 ° C. or higher and 500 ° C. or lower. When further alloying is performed, the temperature during alloying is 550 ° C. or lower. It is preferable.
- alloying temperature exceeds 550 ° C.
- carbide precipitates from untransformed austenite or pearlite is generated in some cases, so that strength and workability or both cannot be obtained, and the powdering property of the plating layer is also low. to degrade.
- the temperature during alloying is less than 450 ° C.
- Coating weight is preferably in a per side 20 g / m 2 or more 150 g / m 2 or less. If the plating adhesion amount is less than 20 g / m 2 , the corrosion resistance is insufficient. On the other hand, if it exceeds 150 g / m 2 g, the corrosion resistance effect is saturated and only the cost is increased.
- the alloying degree (Fe mass% (Fe content)) of the plating layer is preferably in the range of 7 mass% to 15 mass%. If the degree of alloying of the plating layer is less than 7% by mass, unevenness in alloying occurs and the appearance quality deteriorates, or the so-called ⁇ phase is generated in the plating layer and the slidability of the steel sheet deteriorates. On the other hand, when the degree of alloying of the plating layer exceeds 15% by mass, a large amount of hard and brittle ⁇ phase is formed, and the plating adhesion deteriorates.
- the slab obtained by melting the steel having the composition shown in Table 1 is heated to 1200 ° C, the hot-rolled steel sheet finished by hot rolling at 870 ° C is wound up at 650 ° C, and then the hot-rolled steel sheet is pickled. Thereafter, it was cold-rolled at a rolling rate of 65% to obtain a cold-rolled steel sheet having a sheet thickness of 1.2 mm.
- the obtained cold-rolled steel sheet was heat-treated under the conditions shown in Table 2.
- the cooling stop temperature: T in Table 2 is a temperature at which the cooling of the steel sheet is stopped when the steel sheet is cooled from the annealing temperature. Further, some cold-rolled steel sheets were subjected to hot dip galvanizing treatment or alloying hot dip galvanizing treatment.
- the hot dip galvanizing treatment double-side plating was performed so that the plating bath temperature was 463 ° C. and the basis weight (per one side) was 50 g / m 2 .
- the alloying hot dip galvanizing treatment is performed by adjusting the alloying conditions so that the weight per unit area (per one side): 50 g / m 2 and the degree of alloying (Fe mass% (Fe content)) is 9 mass%. Double-sided plating was applied.
- the hot dip galvanizing treatment and the alloying hot galvannealing hot dip galvanizing treatment were performed after cooling to T ° C shown in Table 2 once.
- the amount of retained austenite was determined by measuring the X-ray diffraction intensity after grinding and polishing the steel plate to 1 ⁇ 4 of the plate thickness in the plate thickness direction. For incident X-rays, Co—K ⁇ is used, and from the intensity ratio of each surface of austenite (200), (220), (311) to the diffraction intensity of each surface of ferrite (200), (211), (220). The average value was calculated for the amount of retained austenite.
- the average amount of C in the retained austenite is obtained by calculating the lattice constant from the intensity peaks of the (200), (220), and (311) surfaces of austenite in the X-ray diffraction intensity measurement.
- C amount (mass%) was calculated
- a0 0.3580 + 0.0033 ⁇ [C%] + 0.00095 ⁇ [Mn%] + 0.0056 ⁇ [Al%] + 0.022 ⁇ [N%]
- [X%] Mass% of element X
- mass% of elements other than C was mass% with respect to the whole steel plate.
- TS tensile strength
- T.P. El total elongation
- TS ⁇ T.El product of strength and total elongation
- the stretch flangeability was evaluated in accordance with Japan Iron and Steel Federation standard JFST1001.
- Each steel plate obtained was cut to 100 mm ⁇ 100 mm, a hole with a clearance of 12% of the plate thickness and a diameter of 10 mm was punched out, and then pressed with a wrinkle holding force of 88.2 kN using a die with an inner diameter of 75 mm.
- a 60 ° conical punch was pushed into the hole, the hole diameter at the crack initiation limit was measured, and the critical hole expansion ratio ⁇ (%) was obtained from the equation (1).
- Limit hole expansion ratio ⁇ (%) ⁇ (Df ⁇ D0) / D0 ⁇ ⁇ 100 (1)
- Df is a hole diameter (mm) at the time of crack occurrence
- D0 is an initial hole diameter (mm).
- the product of strength and limit hole expansion rate (TS ⁇ ⁇ ) was calculated using ⁇ measured in this manner, and the balance between strength and stretch flangeability was evaluated. In the present invention, when TS ⁇ ⁇ ⁇ 25000 MPa ⁇ %, the stretch flangeability is good.
- the hardness of the hardest structure in the steel sheet structure was judged by the following method. That is, when martensite is observed as-quenched as a result of structure observation, these martensite as-quenched is measured at 10 points at a load of 0.02N with ultra micro Vickers, and the average value thereof is measured in the steel sheet structure. The hardness of the hardest tissue.
- any of the structures of tempered martensite, upper bainite or lower bainite is the hardest phase in the steel sheet of the present invention. These hardest phases were HV ⁇ 800 in the case of the steel sheet of the present invention.
- Table 3 shows the above evaluation results.
- all the steel plates of the present invention have a tensile strength of 980 MPa or more and TS ⁇ T. Since the value of El satisfies 20000 MPa ⁇ % or more and TS ⁇ ⁇ ⁇ 25000 MPa ⁇ %, it was confirmed that both the high strength and excellent workability, particularly excellent stretch flangeability were obtained.
- sample no. No. 1 has an average cooling rate of up to 550 ° C. that is outside the proper range, so that a desired steel sheet structure cannot be obtained and TS ⁇ ⁇ ⁇ 25000 MPa ⁇ % is satisfied, but tensile strength (TS) ⁇ 980 MPa and TS ⁇ T. EL ⁇ 20000 MPa ⁇ % was not satisfied.
- Sample No. No. 2 is a sample No. 2 because the holding time in the first temperature range is outside the proper range. 5, since the annealing temperature is below A 3 point ° C., Sample No. 6 is the cooling stop temperature: T is outside the first temperature range. Since the holding temperature in the second temperature range is outside the proper range, No.
- TS 11 has a holding time in the second temperature range that is outside the appropriate range, so that a desired steel sheet structure cannot be obtained, and although tensile strength (TS) ⁇ 980 MPa is satisfied, TS ⁇ T. Either EL ⁇ 20000 MPa ⁇ % or TS ⁇ ⁇ ⁇ 25000 MPa ⁇ % was not satisfied. Sample No. 30 to 34, since the component composition is outside the proper range, the desired steel sheet structure cannot be obtained, and the tensile strength (TS) ⁇ 980 MPa, TS ⁇ T. Any one or more of EL ⁇ 20000 MPa ⁇ % and TS ⁇ ⁇ ⁇ 25000 MPa ⁇ % were not satisfied.
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Abstract
Description
特許文献1に記載される成分組成では、鋼板に歪みを付与した際に、高歪域でのTRIP効果を発現する安定した残留オーステナイトの量を確保することが困難であり、曲げ性は得られるものの、塑性不安定が生じるまでの延性が低く、張り出し性に劣る。
本発明の高強度鋼板には、鋼板の表面に溶融亜鉛めっきまたは合金化溶融亜鉛めっきを施した鋼板を含むものとする。
なお、本発明において、加工性に優れるとは、TS×T.ELの値が20000MPa・%以上、かつTS×λの値が25000MPa・%を満足することである。ただし、TSは、引張強さ(MPa)、T.ELは、全伸び(%)、λは、限界穴拡げ率(%)である。
1.質量%で
C:0.17%以上0.73%以下、
Si:3.0%以下、
Mn:0.5%以上3.0%以下、
P:0.1%以下、
S:0.07%以下、
Al:3.0%以下および
N:0.010%以下
を含有し、かつSi+Alが0.7%以上を満足し、残部はFeおよび不可避不純物の組成になり、
鋼板組織として、下部ベイナイトおよび全マルテンサイトの合計量の鋼板組織全体に対する面積率が10%以上90%以下、残留オーステナイト量が5%以上50%以下、上部ベイナイト中のベイニティックフェライトの鋼板組織全体に対する面積率が5%以上であり、前記下部ベイナイトおよび全マルテンサイトの合計量のうち焼入れままのマルテンサイトが75%以下、ポリゴナルフェライトの鋼板組織全体に対する面積率が10%以下(0%を含む)を満足し、かつ前記残留オーステナイト中の平均C量が0.70%以上であって、引張強さが980MPa以上であることを特徴とする高強度鋼板。
Cr:0.05%以上5.0%以下、
V:0.005%以上1.0%以下および
Mo:0.005%以上0.5%以下
のうちから選んだ1種または2種以上の元素を含有することを特徴とする上記1に記載の高強度鋼板。
Ti:0.01%以上0.1%以下および
Nb:0.01%以上0.1%以下
のうちから選んだ1種または2種の元素を含有することを特徴とする上記1または2に記載の高強度鋼板。
B:0.0003%以上0.0050%以下
を含有することを特徴とする上記1乃至3のいずれか1項に記載の高強度鋼板。
Ni:0.05%以上2.0%以下および
Cu:0.05%以上2.0%以下
のうちから選んだ1種または2種の元素を含有することを特徴とする上記1乃至4のいずれか1項に記載の高強度鋼板。
Ca:0.001%以上0.005%以下および
REM:0.001%以上0.005%以下
のうちから選んだ1種または2種の元素を含有することを特徴とする上記1乃至5のいずれか1項に記載の高強度鋼板。
まず、本発明において、鋼板組織を上記のように限定した理由について述べる。以下、面積率は、鋼板組織全体に対する面積率とする。
下部ベイナイトおよびマルテンサイトは、鋼板を高強度化のために必要な組織である。下部ベイナイトおよび全マルテンサイトの合計量の面積率が10%未満では、鋼板の引張強さ(TS)が980MPaを満足しない。一方、下部ベイナイトと全マルテンサイトの合計量の面積率が90%を超えると、上部ベイナイトが少なくなり、結果的にCの濃化した安定な残留オーステナイトが確保できないため、延性等の加工性が低下することが問題となる。従って、下部ベイナイトおよび全マルテンサイトの合計量の面積率は、10%以上90%以下とした。好ましくは、20%以上80%以下の範囲である。より好ましくは、30%以上70%以下の範囲である。
マルテンサイトのうち、焼入れままのマルテンサイトの割合が、鋼板中に存在する下部ベイナイトおよび全マルテンサイトの合計量に対して75%を超えると、引張強さは980MPa以上となるものの、伸びフランジ性に劣る。焼入れままのマルテンサイトは極めて硬質であり、焼入れままのマルテンサイト自体の変形能は極めて低いため、鋼板の加工性とりわけ伸びフランジ性を著しく劣化させる。また、焼入れままのマルテンサイトと上部ベイナイトの硬度差は著しく大きいため、焼入れままのマルテンサイトの量が多いと、焼入れままのマルテンサイトと上部ベイナイトとの界面が多くなり、打ち抜き加工時などに、焼入れままのマルテンサイトと上部ベイナイトの界面に微小なボイドが発生し、打ち抜き加工の後に行う伸びフランジ成形時に、ボイドが連結して亀裂が進展しやすくなることから、伸びフランジ性が劣化する。従って、マルテンサイトのうち焼入れままのマルテンサイトの割合は、鋼板中に存在する下部ベイナイトおよび全マルテンサイトの合計量に対して75%以下とする。好ましくは50%以下である。なお、焼入れままのマルテンサイトは、マルテンサイト中に炭化物が認められない組織で、SEMにより観察することができる。
残留オーステナイトは、加工時にTRIP効果によりマルテンサイト変態し、歪分散能を高めることにより延性を向上させる。
本発明の鋼板では、上部ベイナイト変態を活用して、特に、C濃化量を高めた残留オーステナイトを、上部ベイナイト中に形成せしめる。その結果、加工時に高歪域でもTRIP効果を発現できる残留オーステナイトを得ることができる。このような残留オーステナイトとマルテンサイトを併存させて活用することにより、引張強さ(TS)が980MPa以上の高強度領域でも良好な加工性が得られ、具体的には、TS×T.Elの値を20000MPa以上とすることができ、強度と延性のバランスに優れた鋼板を得ることができる。
ここで、上部ベイナイト中の残留オーステナイトは、上部ベイナイト中のベイニティックフェライトのラス間に形成され、細かく分布するため、組織観察によりその量(面積率)を求めるには高倍率で大量の測定が必要であり、正確に定量することは難しい。しかし、該ベイニティックフェライトのラス間に形成される残留オーステナイトの量は、形成されるベイニティックフェライト量にある程度見合った量である。そこで、発明者らが検討した結果、上部ベイナイト中のベイニティックフェライトの面積率が5%以上で、かつ従来から行われている残留オーステナイト量を測定する手法であるX線回折(XRD)による強度測定、具体的にはフェライトとオーステナイトのX線回折強度比から求められる残留オーステナイト量が5%以上であれば、十分なTRIP効果を得ることができ、引張強さ(TS)が980MPa以上で、TS×T.Elが20000MPa・%以上を達成できることが分かった。なお、従来から行われている残留オーステナイト量の測定手法で得られた残留オーステナイト量は、残留オーステナイトの鋼板組織全体に対する面積率と同等であることを確認している。
残留オーステナイト量が5%未満の場合、十分なTRIP効果が得られない。一方、50%を超えると、TRIP効果発現後に生じる硬質なマルテンサイトが過大となり、靭性の劣化などが問題となる。従って、残留オーステナイトの量は、5%以上50%以下の範囲とする。好ましくは、5%超、より好ましくは10%以上45%以下の範囲である。さらに好ましくは、15%以上40%以下の範囲である。
TRIP効果を活用して優れた加工性を得るためには、引張強さ(TS)が980MPa~2.5GPa級の高強度鋼板においては、残留オーステナイト中のC量が重要である。本発明の鋼板では、上部ベイナイト中のベイニティックフェライトのラス間に形成される残留オーステナイトにCを濃化させる。該ラス間の残留オーステナイト中に濃化されるC量を正確に評価することは困難であるが、発明者らが検討した結果、本発明の鋼板においては、従来行われている残留オーステナイト中の平均C量(残留オーステナイト中のC量の平均)を測定する方法であるX線回折(XRD)での回折ピークのシフト量から求める残留オーステナイト中の平均C量が0.70%以上であれば、優れた加工性が得られることが分かった。
残留オーステナイト中の平均C量が0.70%未満の場合、加工時において低歪域でマルテンサイト変態が生じてしまい、加工性を向上させる高歪域でのTRIP効果が得られない。従って、残留オーステナイト中の平均C量は0.70%以上とする。好ましくは0.90%以上である。一方、残留オーステナイト中の平均C量が2.00%を超えると、残留オーステナイトが過剰に安定となり、加工中にマルテンサイト変態が生じず、TRIP効果が発現しないことにより、延性が低下する。従って、残留オーステナイト中の平均C量は2.00%以下とすることが好ましい。より好ましくは1.50%以下である。
上部ベイナイト変態によるベイニティックフェライトの生成は、未変態オーステナイト中のCを濃化させ、加工時に高歪域でTRIP効果を発現して歪分解能を高める残留オーステナイトを得るために必要である。オーステナイトからベイナイトへの変態は、およそ150~550℃の広い温度範囲にわたって起こり、この温度範囲内で生成するベイナイトには種々のものが存在する。従来技術では、このような種々のベイナイトを単にベイナイトと規定する場合が多かったが、本発明で目標とする加工性を得るためには、ベイナイト組織を明確に規定する必要があることから、上部ベイナイトおよび下部ベイナイトを次のように定義する。
上部ベイナイトは、ラス状のベイニティックフェライトと、ベイニッティクフェライトの間に存在する残留オーステナイトおよび/または炭化物とからなり、ラス状のベイニティックフェライト中に規則正しく並んだ細かな炭化物が存在しないことが特徴である。一方、下部ベイナイトは、ラス状のベイニティックフェライトと、ベイニッティクフェライトの間に存在する残留オーステナイトおよび/または炭化物とからなることは、上部ベイナイトと共通であるが、下部ベイナイトでは、ラス状のベイニティックフェライト中に規則正しく並んだ細かな炭化物が存在することが特徴である。
つまり、上部ベイナイトと下部ベイナイトは、ベイニティックフェライト中における規則正しく並んだ細かな炭化物の有無によって区別される。このようなベイニティックフェライト中における炭化物の生成状態の差は、残留オーステナイト中へのCの濃化に大きな影響を与える。つまり、上部ベイナイトのベイニティックフェライトの面積率が5%未満の場合、ベイナイト変態を進めた場合においても、Cはベイニティックフェライト中に炭化物として生成する量が多くなり、結果的にラス間に存在する残留オーステナイト中へのC濃化量が減少して、加工時に高歪域でTRIP効果を発現する残留オーステナイト量が減少することが問題となる。従って、上部ベイナイト中のベイニティックフェライトの面積率は、鋼板組織全体に対する面積率で5%以上必要である。一方、上部ベイナイトのベイニティックフェライトの鋼板組織全体に対する面積率が85%を超えると、強度の確保が困難となる場合があるため、85%以下とすることが好ましい。
ポリゴナルフェライトの面積率が10%を超えると、引張強さ(TS):980MPa以上を満足することが困難になると同時に、加工時に硬質組織内に混在した軟質なポリゴナルフェライトに歪が集中することにより加工時に容易に亀裂が発生し、結果として所望の加工性を得られない。ここで、ポリゴナルフェライトの面積率が10%以下であれば、ポリゴナルフェライトが存在しても硬質相中に少量のポリゴナルフェライトが孤立分散した状態となり、歪の集中を抑制することができ、加工性の劣化を避けることができる。従って、ポリゴナルフェライトの面積率は10%以下とする。好ましくは5%以下、さらに好ましくは3%以下であり、0%であってもよい。
C:0.17%以上0.73%以下
Cは鋼板の高強度化および安定した残留オーステナイト量を確保するのに必要不可欠な元素であり、マルテンサイト量の確保および室温でオーステナイトを残留させるために必要な元素である。C量が0.17%未満では、鋼板の強度と加工性を確保することが難しい。一方、C量が0.73%を超えると、溶接部および熱影響部の硬化が著しく溶接性が劣化する。従って、C量は0.17%以上0.73%以下の範囲とする。好ましくは、0.20%を超え0.48%以下の範囲であり、さらに好ましくは0.25%以上である。
Siは、固溶強化により鋼の強度向上に寄与する有用な元素である。しかしながら、Si量が3.0%を超えると、ポリゴナルフェライトおよびベイニティックフェライト中への固溶量の増加による加工性、靭性の劣化を招き、また、赤スケール等の発生による表面性状の劣化や、溶融めっきを施す場合には、めっき付着性および密着性の劣化を引き起こす。従って、Si量は3.0%以下とする。好ましくは2.6%以下である。さらに好ましくは、2.2%以下である。
また、Siは、炭化物の生成を抑制し、残留オーステナイトの生成を促進するのに有用な元素であることから、Si量は0.5%以上とすることが好ましいが、炭化物の生成をAlのみで抑制する場合には、Siは添加する必要はなく、Si量は0%であっても良い。
Mnは、鋼の強化に有効な元素である。Mn量が0.5%未満では、焼鈍後の冷却中にベイナイトやマルテンサイトが生成する温度よりも高い温度域で炭化物が析出するため、鋼の強化に寄与する硬質相の量を確保することができない。一方、Mn量が3.0%を超えると、鋳造性の劣化などを引き起こす。従って、Mn量は0.5%以上3.0%以下の範囲とする。好ましくは1.5%以上2.5%以下の範囲とする。
Pは鋼の強化に有用な元素であるが、P量が0.1%を超えると、粒界偏析により脆化することにより耐衝撃性を劣化させ、鋼板に合金化溶融亜鉛めっきを施す場合には、合金化速度を大幅に遅延させる。従って、P量は0.1%以下とする。好ましくは0.05%以下である。なお、P量は、低減することが好ましいが、0.005%未満とするには大幅なコスト増加を引き起こすため、その下限は0.005%程度とすることが好ましい。
Sは、MnSを生成して介在物となり、耐衝撃性の劣化や溶接部のメタルフローに沿った割れの原因となるため、S量を極力低減することが好ましい。しかしながら、S量を過度に低減することは、製造コストの増加を招くため、S量は0.07%以下とする。好ましくは0.05%以下であり、より好ましくは0.01%以下である。なお、Sを0.0005%未満とするには大きな製造コストの増加を伴うため、製造コストの点からはその下限は0.0005%程度である。
Alは、鋼の強化に有用な元素であるとともに、製鋼工程で脱酸剤として添加される有用な元素である。Al量が3.0%を超えると、鋼板中の介在物が多くなり延性を劣化させる。従って、Al量は3.0%以下とする。好ましくは、2.0%以下である。
また、Alは、炭化物の生成を抑制し、残留オーステナイトの生成を促進するのに有用な元素であり、また、脱酸効果を得るために、Al量は、0.001%以上とすることが好ましく、より好ましくは0.005%以上とする。なお、本発明におけるAl量は、脱酸後に鋼板中に含有するAl量とする。
Nは、鋼の耐時効性を最も大きく劣化させる元素であり、極力低減することが好ましい。N量が0.010%を超えると耐時効性の劣化が顕著となるため、N量は0.010%以下とする。なお、Nを0.001%未満とするには大きな製造コストの増加を招くため、製造コストの点からは、その下限は0.001%程度である。
Si+Al≧0.7%
SiおよびAlはともに、上記したように、炭化物の生成を抑制し、残留オーステナイトの生成を促進するのに有用な元素である。炭化物の生成の抑制は、SiまたはAlを単独で含有させても効果があるが、Si量とAl量の合計で0.7%以上を満足する必要がある。なお、上掲式におけるAl量は、脱酸後に鋼板中に含有するAl量とする。
Cr:0.05%以上5.0%以下、V:0.005%以上1.0%以下およびMo:0.005%以上0.5%以下のうちから選ばれる1種または2種以上
Cr、VおよびMoは焼鈍温度からの冷却時にパーライトの生成を抑制する作用を有する元素である。その効果は、Cr:0.05%以上、V:0.005%以上およびMo:0.005%以上で得られる。一方、Cr:5.0%、V:1.0%およびMo:0.5%を超えると、硬質なマルテンサイトの量が過大となり、必要以上に高強度となる。従って、Cr、VおよびMoを含有させる場合には、Cr:0.05%以上5.0%以下、V:0.005%以上1.0%以下およびMo:0.005%以上0.5%以下の範囲とする。
TiおよびNbは鋼の析出強化に有用で、その効果は、それぞれの含有量が0.01%以上で得られる。一方、それぞれの含有量が0.1%を超えると加工性および形状凍結性が低下する。従って、TiおよびNbを含有させる場合は、Ti:0.01%以上0.1%以下およびNb:0.01%以上0.1%以下の範囲とする。
Bはオーステナイト粒界からフェライトが生成・成長することを抑制するのに有用な元素である。その効果は0.0003%以上の含有で得られる。一方、含有量が0.0050%を超えると加工性が低下する。従って、Bを含有させる場合は、B:0.0003%以上0.0050%以下の範囲とする。
NiおよびCuは鋼の強化に有効な元素である。また、鋼板に溶融亜鉛めっきまたは合金化溶融亜鉛めっきを施す場合には、鋼板表層部の内部酸化を促進してめっき密着性を向上させる。これらの効果は、それぞれの含有量が0.05%以上で得られる。一方、それぞれの含有量が2.0%を超えると、鋼板の加工性を低下させる。従って、NiおよびCuを含有させる場合には、Ni:0.05%以上2.0%以下およびCu:0.05%以上2.0%以下の範囲とする。
CaおよびREMは、硫化物の形状を球状化し、伸びフランジ性への硫化物の悪影響を改善するために有用である。その効果は、それぞれの含有量が0.001%以上で得られる。一方、それぞれの含有量が0.005%を超えると、介在物等の増加を招き、表面欠陥および内部欠陥などを引き起こす。従って、CaおよびREMを含有させる場合には、Ca:0.001%以上0.005%以下およびREM:0.001%以上0.005%以下の範囲とする。
上記の好適成分組成に調整した鋼片を製造後、熱間圧延し、ついで冷間圧延を施して冷延鋼板とする。本発明において、これらの処理に特に制限はなく、常法に従って行えば良い。
好適な製造条件は次のとおりである。鋼片を、1000℃以上1300℃以下の温度域に加熱した後、870℃以上950℃以下の温度域で熱間圧延を終了し、得られた熱延鋼板を350℃以上720℃以下の温度域で巻き取る。ついで、熱延鋼板を酸洗後、40%以上90%以下の範囲の圧下率で冷間圧延を行い冷延鋼板とする。
なお、本発明では、鋼板を通常の製鋼、鋳造、熱間圧延、酸洗および冷間圧延の各工程を経て製造する場合を想定しているが、例えば、薄スラブ鋳造やストリップ鋳造などにより熱間圧延工程の一部または全部を省略して製造しても良い。
オーステナイト単相域で15秒以上600秒以下の焼鈍をする。本発明の鋼板は、未変態オーステナイトから、350℃以上490℃以下の範囲の比較的低温域で変態させる上部ベイナイト、下部ベイナイトおよびマルテンサイトを主相とするため、ポリゴナルフェライトが極力少ない方が好ましく、オーステナイト単相域での焼鈍が必要である。焼鈍温度に関しては、オーステナイト単相域であれば特に制限はないが、焼鈍温度が1000℃を超えるとオーステナイト粒の成長が著しく、後の冷却によって生じる構成相の粗大化を引き起こし、靭性などを劣化させる。一方、焼鈍温度がA3点(オーステナイト変態点)未満の場合には、焼鈍段階で既にポリゴナルフェライトが生成しており、冷却中のポリゴナルフェライトの成長を抑制するためには500℃以上の温度域を極めて急速に冷却する必要が生じる。従って、焼鈍温度は、A3点以上とすることが必要であり、1000℃以下とすることが好ましい。
また、焼鈍時間が15秒未満の場合には、オーステナイトへの逆変態が十分に進まない場合や、鋼板中の炭化物が十分に溶解しない場合がある。一方、焼鈍時間が600秒を超えると、多大なエネルギー消費に伴うコスト増を招く。従って、焼鈍時間は15秒以上600秒以下の範囲とする。好ましくは、60秒以上500秒以下の範囲である。ここで、A3点は、
A3点(℃)=910−203×[C%]1/2+44.7×[Si%]−30×[Mn%]
+700×[P%]+130×[Al%]−15.2×[Ni%]
−11×[Cr%]−20×[Cu%]+31.5×[Mo%]
+104×[V%]+400×[Ti%]
によって算出することができる。なお、[X%]は鋼板の成分元素Xの質量%とする。
T℃以上550℃以下の温度域で鋼板が冷却される速度は、該第1保持温度域での保持時間を15秒以上1000秒以下とする以外は特に制限されないが、鋼板が過度に遅い速度で冷却された場合には、未変態オーステナイトから炭化物が生成することにより、所望の組織が得られなくなる可能性が高い。従って、T℃以上550℃以下の温度域において、鋼板は、平均で1℃/s以上の速度で冷却されることが好ましい。
また、第1温度域での保持時間が15秒未満の場合、上部ベイナイト変態量が少なくなり、未変態オーステナイト中へのC濃化量が少なくなることが問題となる。一方、第1温度域での保持時間が1000秒を超える場合、鋼板の最終組織として残留オーステナイトとなる未変態オーステナイトから炭化物が析出してC濃化した安定な残留オーステナイトが得られず、その結果、所望の加工性が得られない。従って、保持時間は15秒以上1000秒以下とする。好ましくは、30秒以上600秒以下の範囲である。
また、保持時間が15秒未満の場合、十分な量の下部ベイナイトを得られず、所望の加工性が得られない。一方、保持時間が1000秒を超えると、第1温度域で生成させた上部ベイナイト中の安定した残留オーステナイトから炭化物が析出し、その結果、所望の加工性が得られない。従って、保持時間は15秒以上1000秒以下の範囲とする。好ましくは、30秒以上600秒以下の範囲である。
また、本発明の製造方法に従い、第2温度域での保持後に引き続き、溶融亜鉛めっき処理あるいは合金化溶融亜鉛めっき処理を施すことができる。
鋼板をめっき浴中に浸入させ、ガスワイピングなどで付着量を調整する。めっき浴中の溶解Al量は、溶融亜鉛めっき処理の場合は0.12%以上0.22%以下の範囲、合金化溶融亜鉛めっき処理の場合は0.08%以上0.18%以下の範囲とすることが好ましい。
処理温度は、溶融亜鉛めっき処理の場合、めっき浴の温度は通常の450℃以上500℃以下の範囲であればよく、さらに合金化処理を施す場合、合金化時の温度は550℃以下とすることが好ましい。合金化温度が550℃を超える場合、未変態オーステナイトから炭化物が析出したり、場合によってはパーライトが生成するため、強度や加工性またはその両方が得られず、また、めっき層のパウダリング性も劣化する。一方、合金化時の温度が450℃未満では合金化が進行しない場合があるため、450℃以上とすることが好ましい。
めっき付着量は片面当たり20g/m2以上150g/m2以下の範囲とすることが好ましい。めっき付着量が20g/m2未満では耐食性が不足し、一方、150g/m2gを超えても耐食効果は飽和し、コストアップを招くだけである。
めっき層の合金化度(Fe質量%(Fe含有量))は7質量%以上15質量%以下の範囲が好ましい。めっき層の合金化度が7質量%未満では、合金化ムラが生じ外観品質が劣化したり、めっき層中にいわゆるζ相が生成され鋼板の摺動性が劣化したりする。一方、めっき層の合金化度が15質量%を超えると、硬質で脆いΓ相が多量に形成され、めっき密着性が劣化する。
また、一部の冷延鋼板については、溶融亜鉛めっき処理あるいは合金化溶融亜鉛めっき処理を施した。ここで、溶融亜鉛めっき処理は、めっき浴温度:463℃、目付け量(片面あたり):50g/m2となるように両面めっきを施した。また、合金化溶融亜鉛めっき処理は、目付け量(片面あたり):50g/m2として合金化度(Fe質量%(Fe含有量))が9質量%となるように合金化条件を調整して両面めっきを施した。なお、溶融亜鉛めっき処理および合金化溶融合金化溶融亜鉛めっき処理は、表2に示すT℃まで一旦冷却した後に行った。
各鋼板から試料を切り出し研磨して、圧延方向に平行な面を走査型電子顕微鏡(SEM)を用いて3000倍で10視野組織観察して、各相の面積率を測定し、各結晶粒の相構造を同定した。
a0=0.3580+0.0033×[C%]+0.00095×[Mn%]+0.0056×[Al%]+0.022×[N%]
ただし、a0:格子定数(nm)、[X%]:元素Xの質量%。なお、C以外の元素の質量%は、鋼板全体に対する質量%とした。
限界穴拡げ率λ(%)={(Df−D0)/D0}×100 ・・・(1)
ただし、Dfは亀裂発生時の穴径(mm)、D0は初期穴径(mm)とする。
このようにして測定したλを用いて強度と限界穴拡げ率の積(TS×λ)を算出して、強度と伸びフランジ性のバランスを評価した。
なお、本発明では、TS×λ≧25000MPa・%の場合、伸びフランジ性を良好とした。
Claims (9)
- 質量%で
C:0.17%以上0.73%以下、
Si:3.0%以下、
Mn:0.5%以上3.0%以下、
P:0.1%以下、
S:0.07%以下、
Al:3.0%以下および
N:0.010%以下
を含有し、かつSi+Alが0.7%以上を満足し、残部はFeおよび不可避不純物の組成になり、
鋼板組織として、下部ベイナイトおよび全マルテンサイトの合計量の鋼板組織全体に対する面積率が10%以上90%以下、残留オーステナイト量が5%以上50%以下、上部ベイナイト中のベイニティックフェライトの鋼板組織全体に対する面積率が5%以上であり、前記下部ベイナイトおよび全マルテンサイトの合計量のうち焼入れままのマルテンサイトが75%以下、ポリゴナルフェライトの鋼板組織全体に対する面積率が10%以下(0%を含む)を満足し、かつ前記残留オーステナイト中の平均C量が0.70%以上であって、引張強さが980MPa以上であることを特徴とする高強度鋼板。 - 前記鋼板がさらに、質量%で、
Cr:0.05%以上5.0%以下、
V:0.005%以上1.0%以下および
Mo:0.005%以上0.5%以下
のうちから選んだ1種または2種以上の元素を含有することを特徴とする請求項1に記載の高強度鋼板。 - 前記鋼板がさらに、質量%で、
Ti:0.01%以上0.1%以下および
Nb:0.01%以上0.1%以下
のうちから選んだ1種または2種の元素を含有することを特徴とする請求項1または2に記載の高強度鋼板。 - 前記鋼板がさらに、質量%で、
B:0.0003%以上0.0050%以下
を含有することを特徴とする請求項1乃至3のいずれか1項に記載の高強度鋼板。 - 前記鋼板がさらに、質量%で、
Ni:0.05%以上2.0%以下および
Cu:0.05%以上2.0%以下
のうちから選んだ1種または2種の元素を含有することを特徴とする請求項1乃至4のいずれか1項に記載の高強度鋼板。 - 前記鋼板がさらに、質量%で、
Ca:0.001%以上0.005%以下および
REM:0.001%以上0.005%以下
のうちから選んだ1種または2種の元素を含有することを特徴とする請求項1乃至5のいずれか1項に記載の高強度鋼板。 - 請求項1乃至6のいずれか1項に記載の鋼板の表面に、溶融亜鉛めっき層または合金化溶融亜鉛めっき層を具えることを特徴とする高強度鋼板。
- 請求項1乃至6のいずれか1項に記載の成分組成になる鋼片を、熱間圧延し、冷間圧延により冷延鋼板とし、ついで該冷延鋼板を、オーステナイト単相域で15秒以上600秒以下の焼鈍をした後、350℃以上490℃以下の第1温度域で定める冷却停止温度:T℃まで冷却するに際し、少なくとも550℃までは平均冷却速度を5℃/s以上に制御して冷却し、その後、該第1温度域で15秒以上1000秒以下保持し、ついで、200℃以上350℃以下の第2温度域で15秒以上1000秒以下保持することを特徴とする高強度鋼板の製造方法。
- 前記冷却停止温度:T℃までの冷却時もしくは前記第1温度域で、溶融亜鉛めっき処理または合金化溶融亜鉛めっき処理を施すことを特徴とする請求項8に記載の高強度鋼板の製造方法。
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KR101341731B1 (ko) | 2013-12-16 |
CA2734978C (en) | 2016-03-29 |
TWI412609B (zh) | 2013-10-21 |
CN102149841B (zh) | 2013-11-20 |
JP2010065273A (ja) | 2010-03-25 |
JP5365112B2 (ja) | 2013-12-11 |
EP2327810B1 (en) | 2019-02-27 |
CN102149841A (zh) | 2011-08-10 |
TW201020329A (en) | 2010-06-01 |
EP2327810A1 (en) | 2011-06-01 |
EP2327810A4 (en) | 2013-11-20 |
CA2734978A1 (en) | 2010-03-18 |
US20110162762A1 (en) | 2011-07-07 |
KR20110042369A (ko) | 2011-04-26 |
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