JP6965956B2 - High-strength steel plate and its manufacturing method - Google Patents

Description

The present invention relates to a high-strength steel plate having excellent workability and a method for producing the same, which is used in an industrial field such as an automobile.

In recent years, from the viewpoint of global environmental conservation, improving the fuel efficiency of automobiles has become an important issue. Therefore, by using a high-strength steel plate as a vehicle body material, there is an active movement to reduce the thickness of vehicle body parts and to reduce the weight of the vehicle body itself.

Generally, in order to increase the strength of a steel sheet, it is necessary to increase the ratio of hard phases such as martensite and bainite to the entire structure of the steel sheet. However, increasing the strength of the steel sheet by increasing the proportion of the hard phase causes a decrease in workability. Therefore, it is desired to develop a steel sheet having both high strength and excellent workability. So far, various composite structure steel sheets have been developed, such as ferrite-martensite double-phase steel (DP steel) and TRIP steel utilizing transformation-induced plasticity of retained austenite.

When the proportion of the hard phase is increased in the composite structure steel sheet, the workability of the steel sheet is strongly influenced by the workability of the hard phase. This is because when the proportion of the hard phase is small and the amount of soft polygonal ferrite is large, the deformability of the polygonal ferrite is dominant over the workability of the steel sheet, and even when the workability of the hard phase is not sufficient. While workability such as ductility was ensured, when the proportion of the hard phase is large, the deformability of the hard phase itself, not the deformability of the polygonal ferrite, directly affects the formability of the steel sheet. Because.

Therefore, in the case of cold-rolled steel sheets, after performing heat treatment to adjust the amount of polygonal ferrite generated in the annealing and subsequent cooling processes, the steel sheets are water-quenched to generate martensite, and the steel sheets are raised again. By warming and keeping it at a high temperature, martensite has been tempered to generate carbides in martensite, which is a hard phase, and the workability of martensite has been improved. However, in the case of the continuous annealing water quenching method in which such water quenching is usually performed, the temperature after quenching is inevitably close to the water temperature, so that most of the untransformed austenite undergoes martensitic transformation. Therefore, it was difficult to utilize retained austenite and other low-temperature metamorphic tissues. Therefore, the improvement of the workability of the hard structure is limited to the effect of tempering martensite, and as a result, the improvement of the workability of the steel sheet is also limited.

Regarding the composite structure steel sheet containing retained austenite, for example, Patent Document 1 defines a predetermined alloy component, and the steel structure is made into a fine and uniform bainite having retained austenite, which is produced by bending workability and bending property. High-strength steel sheets with excellent impact characteristics have been proposed.

Further, Patent Document 2 defines a predetermined alloy component, defines the steel structure as bainite having retained austenite, and specifies the amount of retained austenite in bainite, which is a composite having excellent seizure curability. Structured steel sheets have been proposed.

Further, Patent Document 3 defines a predetermined alloy component, and the steel structure is such that bainite having retained austenite is 95% or more in area ratio, the amount of retained austenite in bainite is 1% or more and 15% or less, and bainite. A composite structure steel sheet having excellent impact resistance, which is manufactured by specifying the hardness (HV) of the above, has been proposed.

Japanese Unexamined Patent Publication No. 4-235253 Japanese Unexamined Patent Publication No. 2004-76114 Japanese Unexamined Patent Publication No. 11-256273

However, the steel sheet described in the above-mentioned prior art document has the following problems.

With the component composition described in Patent Document 1, when strain is applied to a steel sheet, it is difficult to secure a stable amount of retained austenite that exhibits a TRIP effect in a high strain region, and bendability can be obtained. However, the ductility until plastic instability occurs is low, and the overhangability is inferior.

Although the steel sheet described in Patent Document 2 has seizure curability, it has a structure in which bainite or further ferrite is mainly contained and martensite is suppressed as much as possible, so that the steel sheet has a tensile strength (TS) of more than 1180 MPa. It is also difficult to ensure workability when the strength is increased.

The steel sheet described in Patent Document 3 has a main purpose of improving impact resistance, and has bainite having a hardness of HV250 or less as the main phase, and specifically has a structure containing 95% or more of bainite. It is extremely difficult to make the strength more than 1180 MPa.

On the other hand, among automobile parts formed by press working, steel plates used as materials for parts that require particularly high strength, such as door impact beams and bumper reinforcements that suppress deformation in the event of an automobile collision, have a value of 1180 MPa or more, and in the future. Further, it is considered that a tensile strength of 1300 MPa or more is required. Further, structural parts such as members and center pillar inners, which are structural parts having a relatively complicated shape, are desired to have a tensile strength of 980 MPa or more, and further 1180 MPa or more in the future.

The present invention solves the problem that it has been difficult to secure workability due to high strength, and provides a high-strength steel plate having a tensile strength (TS) of 1300 MPa or more and excellent workability and a method for manufacturing the same. The purpose is.

In the present invention, high strength means that the tensile strength (TS) is 1300 MPa or more, and excellent workability means tensile strength (TS) × total elongation (TEl). Means that the value of 18,000 MPa ·% or more and the tensile strength × hole expansion ratio (λ) is 40,000 MPa ·% or more.

In order to solve the above problems, the present inventors have made extensive studies on the composition and microstructure of the steel sheet. As a result, the strength was increased by utilizing the martensite and the lower bainite structure, the C content in the steel plate was increased to 0.10% or more, and heat treatment was performed after hot rolling to locally concentrate Mn. In addition, during continuous annealing after cold rolling, the steel plate annealed in the austenite single-phase region is rapidly cooled to partially transform austenite into maltensite, and then the maltenite is tempered and the lower bainite transformation and retained austenite are stabilized. As a result, it has been found that a high-strength steel plate having an excellent balance between workability, particularly strength and ductility, and a tensile strength of 1300 MPa or more can be obtained.

The present invention is based on the above findings, and its gist structure is as follows.
[1] In terms of mass%, C: 0.10% or more and 0.20% or less, Si: 0.01% or more and 2.50% or less, Mn: 3.5% or more and 6.0% or less, P: 0. A steel plate and a component composition containing 100% or less, S: 0.0500% or less, Al: 0.01% or more and 0.50% or less, and N: 0.010% or less, and the balance is Fe and unavoidable impurities. In terms of area ratio to the entire structure, lower bainite is 30% or more and less than 80%, martensite including tempered martensite is 10% or more and less than 50%, retained austenite is 10% or more and 30% or less, and polygonal ferrite is 10% or less ( (Including 0%), a steel plate structure in which the average Mn amount in the retained austenite is 7% or more in mass% and the average Mn amount in lower bainite, martensite and polygonal ferrite is 4% or less in mass%. A high-strength steel plate having a tensile strength of 1300 MPa or more, a tensile strength × total elongation of 18000 MPa ·% or more, and a tensile strength × hole expansion ratio of 40,000 MPa ·% or more.
[2] Further, the above-mentioned component composition is, in terms of mass%, Cr: 0.005% or more and 1.000% or less, V: 0.005% or more and 1.000% or less, Ni: 0.005% or more and 1. It is characterized by containing one or more selected from 000% or less, Mo: 0.005% or more and 1.000% or less, and Cu: 0.01% or more and 2.00% or less [1]. ] The high-strength steel plate described in.
[3] Further, the above-mentioned component composition is one or two selected from Ti: 0.005% or more and 0.100% or less and Nb: 0.005% or more and 0.100% or less in mass%. The high-strength steel plate according to [1] or [2], which comprises.
[4] The high strength according to any one of [1] to [3], wherein the above-mentioned component composition further contains B: 0.0003% or more and 0.0050% or less in mass%. Steel plate.
[5] The above-mentioned component composition is further one or two selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less in mass%. The high-strength steel plate according to any one of [1] to [4], which comprises.
[6] Further, the component composition is, in mass%, Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.100% or less, W: 0.500% or less, Mg: 0. The high-strength steel plate according to any one of claims [1] to [5], which contains one or more selected from 0050% or less and Zr: 0.1000% or less. ..
[7] After hot rolling a steel piece having the composition according to any one of [1] to [6], 600 s or more and 108000 s in a temperature range of _{(Ac 1} transformation point + 20 ° C.) or more and Ac _{3 transformation point or less.} The steel sheet after the first heat treatment step and the first heat treatment step to be held below is cooled, and the steel sheet is cold-rolled to be a cold-rolled steel sheet, and then the cold-rolled steel sheet is annealed in a single-phase region of austenite for 15 seconds or more. After annealing for less than a second, it is cooled to a first temperature range of (Ms point-100 ° C.) or more and less than Ms point at an average cooling rate of 10 ° C./s or more, and then 250 ° C. or more (Bs point-20 ° C.) or It is characterized by having a second heat treatment step of raising the temperature to a second temperature range of 450 ° C., whichever is lower than the lower temperature, and subsequently holding the temperature in the second temperature range for 15 seconds or more and 1000 seconds or less. A method for manufacturing a strong steel plate.
Here, the Ms point is the martensitic transformation start temperature, and the Bs point is a value obtained by the following formula.
Bs (° C.) = 830-270 x [C%] -90 x [Mn%] -37 x [Ni%] -70 x [Cr%] -83 x [Mo%] However, [X%] is for steel sheets. Let it be the mass% of the component element X.

According to the present invention, it is possible to obtain a high-strength steel plate having a tensile strength of 1300 MPa or more and excellent workability. By applying the high-strength steel plate of the present invention to, for example, an automobile structural member, it is possible to improve fuel efficiency by reducing the weight of the vehicle body, and the industrial utility value of the present invention is extremely large.

FIG. 1 is a diagram for explaining a first heat treatment step after hot rolling and before cold rolling in the method for producing a high-strength steel plate according to the present invention. FIG. 2 is a diagram for explaining a second heat treatment step after cold rolling in the method for producing a high-strength steel plate according to the present invention.

Hereinafter, the present invention will be specifically described.

In the present invention, the reason for limiting the component composition of the steel sheet will be described. In the following,% representing the component composition shall mean mass%.

C: 0.10% or more and 0.20% or less C is an element indispensable for increasing the strength of the steel sheet and ensuring a stable amount of retained austenite, in order to secure the amount of martensite and to retain austenite at room temperature. It is an element necessary for. If the amount of C is less than 0.10%, it is difficult to secure the strength and workability of the steel sheet. On the other hand, when the amount of C exceeds 0.20%, the welded portion and the weld heat-affected zone are significantly hardened and the weldability is significantly deteriorated. Therefore, the amount of C is set in the range of 0.10% or more and 0.20% or less. The amount of C is preferably 0.12% or more, more preferably 0.14% or more. The amount of C is preferably 0.18% or less, more preferably 0.17% or less.

Si: 0.01% or more and 2.50% or less Si is a useful element that contributes to the improvement of steel strength and the suppression of carbides by solid solution strengthening. If the amount of Si is less than 0.01%, the effect of adding the Si becomes poor, so the amount of Si is set to 0.01% or more. The amount of Si is preferably 0.05% or more, more preferably 0.10% or more. On the other hand, if the amount of Si exceeds 2.50%, the embrittlement of the steel sheet reduces the hole expansion rate and makes it difficult to secure the desired stretch flangeability. Therefore, the amount of Si is set to 2.50% or less. The amount of Si is preferably 2.00% or less, more preferably 1.50% or less.

Mn: 3.5% or more and 6.0% or less Mn is an extremely important additive element in the present invention. That is, Mn is an element that stabilizes retained austenite, is effective in ensuring good ductility, and is an element that increases the strength of steel by solid solution strengthening. Further, by enriching Mn in the retained austenite, it is possible to secure a large amount of retained austenite as a volume fraction of 10% or more. Such an action is observed when the amount of Mn is 3.5% or more. On the other hand, if the amount of Mn exceeds 6.0%, an appropriate Mn concentration cannot be obtained. From this point of view, the amount of Mn is 3.5% or more and 6.0% or less. The amount of Mn is preferably 4.0% or more, more preferably 4.2% or more. The amount of Mn is preferably 5.5% or less, more preferably 5.0% or less.

P: 0.100% or less P is an element useful for strengthening steel, but when the amount of P exceeds 0.100%, it becomes brittle due to grain boundary segregation and deteriorates impact resistance, resulting in a steel sheet. When alloying hot dip galvanizing is applied, the alloying rate is significantly delayed. Therefore, the amount of P is set to 0.100% or less. The amount of P is preferably 0.050% or less. The amount of P is more preferably 0.030% or less. The amount of P is preferably reduced, but the amount of P is preferably 0.005% or more because it causes a significant cost increase if it is less than 0.005%. More preferably, it is 0.007% or more.

S: 0.0500% or less S becomes an inclusion such as MnS and causes deterioration of impact resistance and cracking along the metal flow of the welded portion. Therefore, it is preferable to reduce the amount of S as much as possible. However, since 0.0500% is acceptable, the amount of S is set to 0.0500% or less. It is preferably 0.0300% or less, and more preferably 0.0100% or less. If S is less than 0.0003%, a large increase in manufacturing cost is required. Therefore, from the viewpoint of manufacturing cost, the amount of S is preferably 0.0003% or more. The amount of S is more preferably 0.0005% or more, still more preferably 0.0008% or more.

Al: 0.01% or more and 0.50% or less Al is a useful element added as a deoxidizer in the steelmaking process. In order to obtain this effect, it is necessary to add 0.01% or more. On the other hand, if the Al content exceeds 0.50%, the risk of slab cracking during continuous casting increases. Therefore, the amount of Al is set to 0.01% or more and 0.50% or less. The amount of Al is preferably 0.02% or more, more preferably 0.03% or more. The amount of Al is preferably 0.40% or less, more preferably 0.30% or less.

N: 0.010% or less N is an element that most deteriorates the aging resistance of steel, and it is preferable to reduce it as much as possible. If the amount of N exceeds 0.010%, the deterioration of aging resistance becomes remarkable, so the amount of N is set to 0.010% or less. The amount of N is preferably 0.008% or less, more preferably 0.007% or less. In addition, since N is less than 0.001%, a large increase in manufacturing cost is caused. Therefore, from the viewpoint of manufacturing cost, the amount of N is preferably 0.001% or more. The amount of N is more preferably 0.002% or more, still more preferably 0.003% or more.

Further, in the present invention, in addition to the above-mentioned basic components, the following components can be appropriately contained.

Cr, V, Ni, Mo: One or more selected from 0.005% or more and 1.000% or less, Cu: 0.01% or more and 2.00% or less, respectively Cr, V, Ni, Mo And Cu are elements that have the effect of suppressing the formation of pearlite when cooled from the annealing temperature. The effect is obtained with Cr, V, Ni, Mo: 0.005% or more and Cu: 0.01% or more, respectively. On the other hand, if Cr, V, Ni, Mo: each exceeds 1.000% and Cu: 2.00%, the amount of hard martensite becomes excessive, and the required processability cannot be obtained. Therefore, when Cr, V, Ni, Mo and Cu are contained, Cr, V, Ni and Mo: 0.005% or more and 1.000% or less, respectively, and Cu: 0.05% or more and 2.00%. The range is as follows. Cr, V, Ni and Mo are preferably 0.010% or more, more preferably 0.015% or more, respectively. Cu is preferably 0.08% or more, more preferably 0.10% or more. Further, Cr, V, Ni and Mo are preferably 0.800% or less, more preferably 0.500% or less, respectively. Cu is preferably 1.00% or less, more preferably 0.80% or less.

One or two Ti and Nb selected from Ti: 0.005% or more and 0.100% or less and Nb: 0.005% or more and 0.100% or less are useful for precipitation strengthening of steel, and their effects are effective. , Each content is obtained at 0.005% or more. On the other hand, if the content of each exceeds 0.100%, the processability and shape freezing property are lowered. Therefore, when Ti and Nb are contained, the range is Ti: 0.005% or more and 0.100% or less and Nb: 0.005% or more and 0.100% or less. When Ti and Nb are contained, Ti: 0.008% or more, Nb: 0.008% or more, more preferably Ti: 0.010% or more, Nb: 0.010% or more. When Ti and Nb are contained, Ti: 0.080% or less, Nb: 0.080% or less, more preferably Ti: 0.060% or less, Nb: 0.060% or less.

B: 0.0003% or more and 0.0050% or less B is an element useful for suppressing the formation and growth of polygonal ferrite from the austenite grain boundary. The effect is obtained with a content of 0.0003% or more. On the other hand, if the B content exceeds 0.0050%, the workability is lowered. Therefore, when B is contained, the range is 0.0003% or more and 0.0050% or less. When B is contained, it is preferably 0.0004% or more, more preferably 0.0005% or more. When B is contained, it is preferably 0.0040% or less, more preferably 0.0030% or less.

In the steel sheet of the present invention, components other than the above are Fe and unavoidable impurities. However, the inclusion of components other than the above is not refused as long as the effects of the present invention are not impaired. The following can be considered as examples of such components.

One or two Ca and REM selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less are all processable by controlling the morphology of sulfide. It is an effective element to improve. In order to obtain such an effect, the content of at least one element selected from Ca and REM needs to be 0.001% or more. The content of at least one element selected from Ca and REM is preferably 0.0015% or more, more preferably 0.0020% or more. On the other hand, if the respective contents of Ca and REM exceed 0.005%, the cleanliness of the steel may be adversely affected. Therefore, the contents of Ca and REM are set to 0.001 to 0.005%, respectively. The content of at least one element selected from Ca and REM is preferably 0.0045% or less, more preferably 0.0040% or less.
Further, in addition to the above-mentioned component composition, the high-strength steel plate of the present invention further has Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.100% or less, W: in mass%. It is preferable to contain one or more selected from 0.500% or less, Mg: 0.0050% or less, and Zr: 0.1000% or less.

Sb is useful for suppressing decarburization of the steel sheet surface during annealing. If the Sb content exceeds 0.200%, the workability is lowered. Therefore, when Sb is contained, it is set to 0.200% or less. When Sb is contained, it is preferably 0.002% or more, more preferably 0.010% or more. When Sb is contained, it is preferably 0.180% or less, more preferably 0.160% or less.
Sn is useful for suppressing decarburization of the steel sheet surface during annealing. If the Sn content exceeds 0.200%, the workability is lowered. Therefore, when Sn is contained, it is set to 0.200% or less. When Sn is contained, it is preferably 0.002% or more, more preferably 0.010% or more. When Sn is contained, it is preferably 0.180% or less, more preferably 0.160% or less.
Ta forms fine carbides, nitrides or carbonitrides during hot rolling or annealing and is useful for precipitation strengthening of steel. If the Ta content exceeds 0.100%, the workability is lowered. Therefore, when Ta is contained, it should be 0.100% or less. When Ta is contained, it is preferably 0.0010% or more, more preferably 0.010% or more. When Ta is contained, it is preferably 0.080% or less, more preferably 0.060% or less.
W forms fine carbides, nitrides or carbonitrides during hot rolling or annealing and is useful for precipitation strengthening of steel. If the W content exceeds 0.500%, the workability is lowered. Therefore, when W is contained, it is set to 0.500% or less. When W is contained, it is preferably 0.0050% or more, more preferably 0.050% or more. When W is contained, it is preferably 0.400% or less, more preferably 0.300% or less.
Mg is an element effective for improving workability by controlling the morphology of inclusions. On the other hand, if the Mg content exceeds 0.0050%, the cleanliness of the steel may be adversely affected. Therefore, when Mg is contained, the Mg content is set to 0.0050% or less. When Mg is contained, it is preferably 0.0005% or more, more preferably 0.0010% or more. When Mg is contained, it is preferably 0.0040% or less, more preferably 0.0030% or less.
Zr is an element effective for improving workability by controlling the morphology of inclusions. On the other hand, if the Zr content exceeds 0.1000%, the cleanliness of the steel may be adversely affected. Therefore, when Zr is contained, the Zr content is set to 0.1000% or less. When Zr is contained, it is preferably 0.0005% or more, more preferably 0.0010% or more. When Zr is contained, it is preferably 0.0900% or less, more preferably 0.0800% or less.

Next, the reason why the steel sheet structure is limited as described above in the present invention will be described. Hereinafter, the area ratio is defined as the area ratio with respect to the entire steel sheet structure in the rolling direction cross section and the plate thickness 1/4 surface position.

The area ratio of retained austenite in the cold-rolled base metal is 10% or more and 30% or less, the amount of Mn in retained austenite is 7% or more by mass%, and the amount of Mn in polygonal ferrite is 4% or less by mass% (0%). including)
In the present invention, it is important to locally concentrate Mn before cold rolling, and when the concentration is 7% or more, excellent workability can be obtained. In order to locally concentrate, the amount of Mn in the polygonal ferrite other than the concentrated portion needs to be 4% or less (including 0%) in mass%. Since Mn is concentrated in austenite, the area ratio of retained austenite needs to be 10% or more. On the other hand, if it exceeds 30%, it is not sufficiently concentrated, so the area ratio of retained austenite is set to 30% or less.

As will be described later, after hot rolling, annealing is performed in the two-phase region of ferrite and austenite for a long time (for example, holding 600s or more and 108000s or less in a temperature region of _{(Ac 1} transformation point + 20 ° C.) or more and Ac _{3 transformation point or less).} Mn is concentrated in austenite, and after cooling, a cold-rolled base material having a Mn-concentrated portion locally can be obtained. Further, since the diffusion coefficient of Mn is small, the Mn-enriched portion locally formed after the hot-rolling annealing is not eliminated in the cold-rolling annealing step, and remains as the Mn-enriched portion.

Since Mn is an element that lowers the Ms point, the Ms point in the Mn-enriched portion is lower than in the case where Mn is not enriched, but conversely, the Ms point other than the Mn-enriched portion where the amount of Mn is small is high. Therefore, when the cooling is stopped in the cold rolling annealing step, the parts other than the Mn-enriched portion are cooled to the Ms point or less when the cooling is stopped. Therefore, martensite is partially generated, and lower bainite is generated in the subsequent reheating. On the other hand, the Mn-enriched portion having a low Ms point is not cooled to below the Ms point when cooling is stopped and does not undergo martensitic transformation, so that retained austenite can be finally secured and excellent ductility can be obtained.

However, if the cold-rolled base metal is not annealed for a long time in the two-phase region of ferrite and austenite after hot rolling, the structure of the cold-rolled base metal is not controlled and there is no Mn-enriched portion with a low Ms point. At the time of stopping, the temperature is cooled to a temperature lower than the Ms point, martensitic transformation proceeds, it becomes difficult to secure retained austenite, and excellent ductility cannot be obtained. Having a Mn-enriched portion with a low Ms point is important for obtaining a cold-rolled annealed material having excellent ductility.

The above is the regulation regarding the cold-rolled base metal, and thereafter, it is the regulation regarding the steel sheet after final annealing.

Area ratio of lower bainite: 30% or more and less than 80% The formation of bainitic ferrite by bainite transformation enriches C in untransformed austenite, and exhibits a TRIP effect in a high strain region during processing to enhance strain resolution. Required to obtain retained austenite. The transformation from austenite to bainite occurs over a wide temperature range of approximately 150-550 ° C., and there are various bainites produced within this temperature range. In the prior art, such various bainites were often simply defined as bainite, but in order to obtain the strength and workability targeted by the present invention, it is necessary to clearly define the bainite structure. , Upper bainite and lower bainite are defined as follows.

The upper bainite consists of lath-like bainite ferrite and retained austenite and / or carbides present between the bainitic ferrites, and there are no regularly arranged fine carbides in the lath-like bainite ferrite. Is a feature. On the other hand, it is common with the upper bainite that the lower bainite is composed of the lath-shaped bainite ferrite and the retained austenite and carbides existing between the bainitic ferrites, but in the lower bainite, the lath-shaped bay is used. It is characterized by the presence of regularly arranged fine carbides in nitic ferrite. That is, upper bainite and lower bainite are distinguished by the presence or absence of regularly arranged fine carbides in bainitic ferrite.

Such a difference in the formation state of carbides in bainitic ferrite has a great influence on the strength of the steel sheet. The upper bainite is softer than the lower bainite, and it is necessary to set the area ratio of the lower bainite to 30% or more in order to obtain the tensile strength targeted in the present invention. The area ratio of the lower bainite is preferably 35% or more, more preferably 40% or more. On the other hand, if the area ratio of the lower bainite is 80% or more, it becomes impossible to obtain retained austenite sufficient for workability, so the area ratio is set to less than 80%. The area ratio of the lower bainite is preferably less than 75%, more preferably less than 70%.

Area ratio of martensite including tempered martensite: 10% or more and less than 50% Martensite is a hard phase and increases the strength of the steel sheet. It also promotes bainite transformation by generating martensite before bainite transformation. If the area ratio of martensite is less than 10%, the bainite transformation cannot be sufficiently promoted, and the bainite area ratio described later cannot be achieved. On the other hand, if the area ratio of martensite exceeds 50%, the bainite structure is reduced and a stable amount of retained austenite cannot be secured, so that there is a problem that processability such as ductility is lowered. Therefore, the area ratio of martensite is 10% or more and less than 50%. The area ratio of martensite is preferably 15% or more, more preferably 20% or more. The area ratio of martensite is preferably less than 40%, more preferably less than 35%. In addition, martensite includes tempered martensite.

Martensite needs to be clearly distinguished from the above-mentioned upper bainite. Martensite can be identified by tissue observation, unquenched martensite does not contain carbides in the tissue, and tempered martensite has carbides with random growth directions in the tissue.

Percentage of tempered martensite out of all martensite: 80% or more As-quenched martensite is extremely hard, has low deformability and is inferior in toughness. As a result, excellent ductility and stretch brittleness may not be obtained. Therefore, due to the significant improvement in the deformability of martensite itself by tempering the martensite as it is hardened, brittle fracture does not occur at the time of applying strain, and the entire structural composition of the present invention including the tempered martensite is realized. By TS × T. El may be 18,000 MPa ·% or more, and TS × λ may be 40,000 MPa ·% or more. Therefore, the proportion of tempered martensite in the martensite is preferably 80% or more of the total martensite area ratio present in the steel sheet. More preferably, the proportion of tempered martensite is 90% or more of the total martensite area ratio.

The tempered martensite is observed as a structure in which fine carbides are precipitated in the martensite by observation with a scanning electron microscope (SEM), and the as-quenched martensite in which such carbides are not observed inside the martensite. It can be clearly distinguished from the site.

Amount of retained austenite: 10% or more and 30% or less Residual austenite undergoes martensitic transformation due to the TRIP effect during processing, and hard martensite containing high C promotes high strength and at the same time increases ductility by increasing strain dispersibility. Improve.

In the steel sheet of the present invention, after partial martensitic transformation, for example, lower bainite transformation in which the formation of carbides is suppressed is utilized to form retained austenite in which the amount of carbon enrichment is increased. As a result, it is possible to obtain retained austenite capable of exhibiting the TRIP effect even in a high strain region during processing.

By coexisting and utilizing such retained austenite, lower bainite and martensite, good workability can be obtained even in a high strength region where the tensile strength (TS) is 1300 MPa or more. Specifically, TS × T .. The El value can be 18,000 MPa ·% or more, and the TS × λ value can be 40,000 MPa ·% or more, so that a steel sheet having an extremely excellent balance between strength and workability can be obtained.

Here, since retained austenite is distributed in a state surrounded by martensite and lower bainite, it is difficult to accurately quantify the amount (area ratio) by observing the structure, but the amount of retained austenite that has been conventionally performed Intensity measurement by X-ray diffraction (XRD), which is a method for measuring, specifically, if the amount of retained austenite obtained from the X-ray diffraction intensity ratio of ferrite and austenite is 10% or more, a sufficient TRIP effect can be obtained. The tensile strength (TS) is 1300 MPa or more, and TS × T. It has been confirmed that El can achieve 18,000 MPa ·% or more. It has been confirmed that the amount of retained austenite obtained by the conventional method for measuring the amount of retained austenite is equivalent to the area ratio of retained austenite to the entire steel sheet structure.

If the amount of retained austenite is less than 10%, a sufficient TRIP effect cannot be obtained. On the other hand, if it exceeds 30%, the hard martensite generated after the manifestation of the TRIP effect becomes excessive, and deterioration of toughness and stretch flangeability becomes a problem. Therefore, the amount of retained austenite should be in the range of 10% or more and 30% or less. Preferably, the amount of retained austenite is 14% or more. Preferably, the amount of retained austenite is 25% or less. A more preferred amount of retained austenite is 18% or more. A more preferred amount of retained austenite is 22% or less.

Area ratio of polygonal ferrite: 10% or less (including 0%)
If the area ratio of the polygonal ferrite exceeds 10%, it becomes difficult to satisfy the tensile strength (TS) of 1300 MPa or more, and at the same time, the strain is concentrated on the soft polygonal ferrite mixed in the hard structure during processing. As a result, cracks are easily generated during processing, and as a result, the desired processability cannot be obtained. Here, if the area ratio of the polygonal ferrite is 10% or less, even if the polygonal ferrite is present, a small amount of the polygonal ferrite is isolated and dispersed in the hard phase, and the concentration of strain can be suppressed. , Deterioration of workability can be avoided. Therefore, the area ratio of polygonal ferrite is set to 10% or less. The area ratio of the polygonal ferrite is preferably 5% or less, more preferably 3% or less, and may be 0%.

The method for measuring the area ratio of each tissue is the method described in Examples described later.

Average Mn amount in retained austenite: 7% or more in mass% In order to obtain excellent workability by utilizing the TRIP effect, in high-strength steel sheets with a tensile strength (TS) of 1300 MPa class or more, in retained austenite The amount of Mn is important. As a result of examination by the inventors, it was found that in the steel sheet of the present invention, even more excellent workability can be obtained when the average amount of Mn in retained austenite is 7% or more. If the average amount of Mn in the retained austenite is less than 7%, martensitic transformation may occur in the low strain region during processing, and the TRIP effect in the high strain region that improves workability may not be sufficiently obtained. .. Therefore, the average amount of Mn in retained austenite is 7% or more. The average amount of Mn in the retained austenite is preferably 8% or more, more preferably 10% or more.

Average Mn amount in lower bainite, martensite, and polygonal ferrite: 4% or less in mass% As described above, in order to make the average Mn amount in retained austenite 7% or more in mass%, the portion other than retained austenite. That is, it is necessary to reduce the average amount of Mn in the lower bainite, martensite, and polygonal ferrite. By setting the average Mn amount in the lower bainite, martensite, and polygonal ferrite to 4% or less in mass%, the average Mn amount in the retained austenite can be sufficiently increased. Therefore, the average amount of Mn in the lower bainite, martensite, and polygonal ferrite is 4% or less in mass%. It is preferably 3.8% or less, more preferably 3.5% or less. The average amount of Mn in the lower bainite, martensite, and polygonal ferrite is preferably 2% or more, more preferably 2.2% or more.

The average amount of Mn in the lower bainite, martensite, and polygonal ferrite can be determined as follows. That is, EPMA (Electron Probe Micro Analyzer) is used to quantify the distribution state of Mn in each structure of the cross section in the rolling direction at the plate thickness 1/4 position. Next, 30 of each of the above 3 tissues are sampled, the amount of Mn thereof is analyzed, and the amount of Mn of each tissue (3 tissues x 30 = 90 tissues) obtained from the analysis result is averaged (simple average). By doing so, it can be obtained.

Next, the method for producing the high-strength steel sheet of the present invention will be described.

After producing the steel piece adjusted to the above-mentioned suitable composition, it is hot-rolled to obtain a hot-rolled steel sheet. In the present invention, these treatments are not particularly limited and may be carried out according to a conventional method, but suitable production conditions are as follows.

After heating the steel piece to a temperature range of 1000 ° C. or higher and 1300 ° C. or lower, hot rolling is completed 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. Wind up in the area. Then, the hot-rolled steel sheet is pickled.

The obtained hot-rolled steel sheet is subjected to the following heat treatment.

Performing the first heat treatment step of long-term annealing, for example, holding 600 s or more and 108,000 s or less in a two-phase region of ferrite and austenite, for example, in _{a temperature region of (Ac 1} transformation point + 20 ° C.) or more and Ac _{3 transformation point or less.} Therefore, the cold-rolled steel sheet of the present invention can be obtained, and the heat treatment for holding the steel sheet is extremely important in the present invention.

That is, _{when the austenite is held in a temperature range lower than (Ac 1} transformation point + 20 ° C.) or a temperature range _{exceeding the Ac 3} transformation point, or when the holding time is less than 600 s, the concentration of Mn in austenite does not proceed. In addition, it becomes difficult to secure a sufficient volume ratio of retained austenite after the final annealing, and the ductility is lowered. On the other hand, when the holding time exceeds 108,000 s, the concentration of Mn in austenite is saturated, which not only reduces the effect on ductility after final annealing, but also causes an increase in cost.

Therefore, the holding temperature in the first heat treatment step of hot rolling plate annealing is (Ac ₁ transformation point + 20 ° C.) or more and Ac ₃ transformation point or less, preferably (Ac ₁ transformation point + 25 ° C.) or more (Ac ₁ transformation point). + 100 ° C. or lower), more preferably (Ac ₁ transformation point + 30 ° C.) or higher (Ac ₁ transformation point + 80 ° C. or lower). The holding time is 600 s or more and 108,000 s or less, preferably 1000 s or more and 80,000 s or less, and more preferably 1200 s or more and 55,000 s or less.

The heat treatment method in the first heat treatment step may be any of batch annealing such as continuous annealing and box annealing. Further, after the first heat treatment step, the cooling is performed to room temperature, but the cooling method and cooling rate are not particularly specified, and any of furnace cooling in batch annealing, gas jet cooling in air cooling and continuous annealing, mist cooling, water cooling, etc. It doesn't matter if it is cooled.

In cold rolling, the rolling reduction is preferably 30% or more. By cold rolling at a rolling reduction of 30% or more, austenite is finely generated during heat treatment, and finally fine retained austenite and martensite are obtained, which not only improves the balance between strength and ductility, but also improves the balance between strength and ductility. Bendability and malleability (hole expandability) may also be improved.

The upper limit of the rolling reduction in cold rolling is not particularly limited, but it is preferably 85% or less from the viewpoint of the load of cold rolling. More preferably, it is 75% or less.

After cold rolling, a second heat treatment step is performed. First, the cold-rolled steel sheet is annealed in the austenite single-phase region for 15 seconds or more and 1000 seconds or less. The steel sheet of the present invention has a low-temperature transformed phase obtained by transforming from untransformed austenite such as martensite as the main phase, and it is preferable that the amount of polygonal ferrite is as small as possible. Therefore, annealing in the austenite single phase region is difficult. is necessary. The annealing temperature is not particularly limited as long as it is in the austenite single-phase region, but when the annealing temperature exceeds 1000 ° C, the growth of austenite grains is remarkable, causing coarsening of the constituent phase caused by subsequent cooling and deterioration of toughness and the like. Let me. Therefore, the annealing temperature is _{preferably Ac 3} points (austenite transformation point) or more, and more preferably (Ac ₃ transformation point + 15 ° C.) or more. The annealing temperature is preferably 1000 ° C. or lower. More preferably, it is 950 ° C. or lower.
Here, Ac ₁ can be calculated by the following equation. [X%] is the mass% of the component element X of the steel sheet.
Ac ₁ point (° C) = 723 + 29 × [Si%] -11 × [Mn%] -17 × [Ni%] + 17 × [Cr%]
Ac ₃ can be calculated by the following equation. [X%] is the mass% of the component element X of the steel sheet.
Ac ₃ points (° C) = 937.2-436.5 x [C%] +56 x [Si%] -19.7 x [Mn%] -26.6 x [Ni%] -4.9 x [Cr %] -16.3 x [Cu%] + 38.1 x [Mo%] +124.8 x [V%] +136.3 x [Ti%] -19.1 x [Nb%] + 198.4 x [Al %] + 3315 x [B%]
Further, when the annealing time in the second heat treatment step is less than 15 seconds, the reverse transformation to austenite may not proceed sufficiently, or the carbides in the steel sheet may not be sufficiently dissolved. On the other hand, if the annealing time in the second heat treatment step exceeds 1000 seconds, the cost increases due to a large amount of energy consumption. Therefore, the annealing time in the second heat treatment step is in the range of 15 seconds or more and 1000 seconds or less. The annealing time in the second heat treatment step is preferably 30 seconds or longer, more preferably 60 seconds or longer. The annealing time in the second heat treatment step is preferably 800 seconds or less, more preferably 600 seconds or less.

In the second heat treatment step, the annealed cold-rolled steel sheet is cooled to a first temperature range of (Ms point −100 ° C.) or more and less than Ms point by controlling the average cooling rate to 10 ° C./s or more. This cooling is to transform a part of austenite into martensitic transformation by cooling to less than the Ms point. Here, if the lower limit of the first temperature range is less than (Ms-100 ° C.), the amount of untransformed austenite converted to martensite becomes excessive at this point, and excellent strength and workability cannot be achieved at the same time. On the other hand, when the upper limit of the first temperature range is Ms or more, an appropriate amount of martensite cannot be secured. Therefore, the range of the first temperature range is set to be (Ms point-100 ° C.) or more and less than the Ms point. It is preferably (Ms point −80 ° C.) or more and less than Ms point, and more preferably (Ms point −50 ° C.) or more and less than Ms point.

Further, in the second heat treatment step, when the average cooling rate from the annealing temperature to the first temperature range is less than 10 ° C./s, excessive formation and growth of polygonal ferrite and precipitation of pearlite and the like occur, which is desired. Steel plate structure cannot be obtained. Therefore, the average cooling rate from the annealing temperature to the first temperature range is set to 10 ° C./s or more. The average cooling rate from the annealing temperature to the first temperature range is preferably 12 ° C./s or more, and more preferably 15 ° C./ or more. The upper limit of the average cooling rate from the annealing temperature to the first temperature range is not particularly limited as long as the cooling shutdown temperature does not vary. The above-mentioned Ms point is preferably determined by thermal expansion measurement during cooling by a four-master test or the like or actual measurement by electrical resistance measurement, but it can also be obtained by an approximate expression as shown in the following equation, for example. Ms is an empirically determined approximation.
Ms (° C.) = 550-35 x [Mn%] -13 x [Si%] -10 x [Cr%] -18 x [Ni%] -12 x [Mo%] -600 x (1-exp (-exp) 0.96 x [C%])))
However, [X%] is the mass% of the component element X of the steel sheet.

The steel sheet cooled to the first temperature range is heated to the second temperature range of 250 ° C. or higher, (Bs point -20 ° C.) or 450 ° C., whichever is lower, and is heated to the second temperature range for 15 seconds in the second temperature range. It is held for a time of 1000 seconds or less.
Bs indicates the bainite transformation start temperature, and is preferably determined by thermal expansion measurement during cooling by a formaster test or the like or actual measurement by electrical resistance measurement, but it can also be determined by an approximate expression as shown in the following equation, for example. Bs is an empirically determined approximate value.
Bs (° C.) = 830-270 x [C%] -90 x [Mn%] -37 x [Ni%] -70 x [Cr%] -83 x [Mo%]
However, [X%] is the mass% of the component element X of the steel sheet.

In the second temperature range, martensite generated by cooling from the annealing temperature to the first temperature range is tempered, untransformed austenite is transformed into lower bainite, and solid solution C is concentrated in austenite to stabilize austenite. Advance the conversion. When the upper limit of the second temperature range exceeds (Bs-20 ° C.) or 450 ° C., bainite transformation is suppressed. On the other hand, when the lower limit of the second temperature range is less than 250 ° C., the diffusion rate of the solid solution C is remarkably lowered, and the amount of C concentration in austenite is reduced, so that the required C concentration in the retained austenite can be obtained. No. Therefore, the range of the second temperature range is 250 ° C. or higher, (Bs-20 ° C.) or 450 ° C., whichever is lower. The range of the second temperature range is preferably 320 ° C. or higher, whichever is lower than (Bs-50 ° C.) or 420 ° C., whichever is lower.

Further, when the holding time in the second temperature range is less than 15 seconds, the tempering of martensite and the transformation of the lower bainite become insufficient, and the desired steel sheet structure cannot be obtained, resulting in the workability of the obtained steel sheet. It may not be possible to secure a sufficient amount of water, so the holding time in this second temperature range needs to be 15 seconds or more. On the other hand, in the present invention, a holding time in the second temperature range of 1000 seconds is sufficient due to the bainite transformation promoting effect of martensite generated in the first temperature range.

Normally, when the amount of alloy components such as C, Cr, and Mn increases as in the steel of the present invention, the bainite transformation is delayed, but when martensite and untransformed austenite coexist as in the present invention, the bainite transformation rate is remarkably high. There have been some reports on this, and the inventors have also found out about the steel of the present invention. On the other hand, when the holding time in the second temperature range exceeds 1000 seconds, carbides are precipitated from untransformed austenite which becomes retained austenite as the final structure of the steel sheet, and stable C-concentrated retained austenite cannot be obtained. As a result, the desired strength and / or ductility may not be obtained. Therefore, the holding time in the second temperature range is set to 15 seconds or more and 1000 seconds or less. The holding time in the second temperature range is preferably 50 seconds or longer, more preferably 100 seconds or longer. The holding time in the second temperature range is preferably 700 seconds or less, more preferably 600 seconds or less.

In the series of heat treatments (first and second heat treatment steps) in the present invention, the holding temperature does not have to be constant as long as it is within the above-mentioned predetermined temperature range, and fluctuates within the predetermined temperature range. Does not impair the gist of the present invention. The same applies to the cooling rate. Further, the steel sheet may be heat-treated by any equipment as long as the heat history is satisfied. Further, it is also included in the scope of the present invention to perform temper rolling on the surface of the steel sheet for shape correction after the heat treatment.

Further, the high-strength hot-dip galvanized steel sheet according to the present invention can be plated to obtain a high-strength hot-dip galvanized steel sheet, or further alloyed to obtain a high-strength hot-dip galvanized steel sheet.

Examples of the present invention are shown below. The present invention is not limited to the following examples.

A steel piece obtained by melting steel having the composition shown in Table 1 is heated to 1250 ° C., and a hot-rolled steel sheet finished and hot-rolled at 870 ° C. is wound at 550 ° C., and then the hot-rolled steel sheet is pickled. After that, heat treatment was performed under the conditions shown in Table 2-1 to perform the first heat treatment step, and then cold rolling was performed at a reduction rate of 50% to obtain a cold-rolled steel sheet having a plate thickness of 1.2 mm. The obtained cold-rolled steel sheet was heat-treated in the second heat treatment step under the conditions shown in Table 2-1. The cooling stop temperature in Table 2-1: T1 is a temperature at which the cooling of the steel sheet in the first temperature range is stopped when the steel sheet is cooled from the annealing temperature. Further, the obtained steel sheet was subjected to temper rolling with a reduction ratio of 0.1%.

The heat treatment in the first heat treatment step after the hot rolling and before the cold rolling generally takes the temperature pattern as shown in FIG. 1, and the heat treatment in the second heat treatment step after the cold rolling is generally shown in FIG. Take a temperature pattern as shown in.

The area ratio of each phase to the entire structure was determined by observing the cross section in the rolling direction and the position of 1/4 of the plate thickness with an optical microscope. Using a cross-sectional tissue photograph at a magnification of 1000 times, the occupied area existing in a square region of 100 μm × 100 μm square arbitrarily set was determined by image analysis. The observation was carried out at N = 5 (5 observation fields). In addition, when observing the tissue, 3 vol. % Piclar and 3 vol. Etched with a mixture of% sodium pyrosulfite.

Assuming that the black region observed in the above microstructure observation is ferrite (polygonal ferrite) or lower bainite, the total area ratio of ferrite and lower bainite was determined. Further, assuming that the remaining region other than the black region is tempered martensite, martensite and retained austenite, the total area ratio of tempered martensite, martensite and retained austenite was determined.

The lower bainite is an aggregate of lath-like ferrite and has a structure containing Fe-based carbides. The lower bainite and the polygonal ferrite are distinguished by using SEM (scanning electron microscope) and TEM (transmission electron microscope). , The area ratio was calculated.

The amount of retained austenite was determined by X-ray diffraction using Mo Kα rays. That is, using a test piece whose measurement surface is a surface near 1/4 of the thickness of the steel sheet, retained austenite is determined from the peak intensities of the (211) and (220) surfaces of austenite and the (200) and (220) surfaces of ferrite. The amount (volume fraction) of was calculated and used as the area ratio. To distinguish between tempered martensite and martensite, a 50 μm × 50 μm square square arbitrarily set by image analysis using a microstructure photograph of a cross section in the rolling direction at a magnification of 1000 to 3000 times observed with a scanning electron microscope (SEM). The occupied area existing in the area was calculated. The observation was carried out at N = 5 (5 observation fields). When observing the structure, etching is performed with nital, and martensite is defined as lumpy and smooth surface on the SEM photograph, and tempered martensite is defined as lumpy and charcoal on the surface. rice field.

The amount of Mn in the retained austenite can be determined as follows. That is, EPMA (Electron Probe Micro Analyzer) is used to quantify the distribution of Mn in each phase of the rolling direction cross section at the plate thickness 1/4 position, and then 30 retained austenite grains. It can be obtained by analyzing the amount of Mn in the above and averaging the amount of Mn of each retained austenite grain obtained from the analysis result.

The average amount of Mn in the lower bainite, martensite, and polygonal ferrite was determined by the method described above in the embodiment for carrying out the invention.

Various characteristics of the obtained steel sheet were evaluated by the following methods.

The tensile test was carried out in accordance with JIS Z 2241 using a JIS No. 5 test piece (JIS Z 2201) in which the plate width direction of the steel sheet was the longitudinal direction. TS (tensile strength) and T.El (total elongation) were measured, and the product of tensile strength and total elongation (TS × T.El) was calculated to evaluate the balance between strength and workability (ductility). In the present invention, the case where TS × T.El ≧ 18000 (MPa ·%) was considered good. Further, a test piece of 100 mm × 100 mm was collected, and a hole expansion test was performed three times in accordance with JFST 1001 (Iron and Steel Federation standard) to obtain an average hole expansion rate (%), and the elongation flange property was evaluated. The product of the tensile strength and the hole expansion rate (TS × λ) was calculated, and the balance between the strength and the workability (elongation flange property) was evaluated. In the present invention, the case of TS × λ ≧ 40,000 (MPa ·%) was regarded as good.

The above evaluation results are shown in Table 2-2.
As is clear from Table 2-2, all the steel sheets produced by the method of the present invention have a tensile strength of 1300 MPa or more, a TS × T.El value of 18000 MPa ·% or more, and a TS × λ value. Was 40,000 MPa ·% or more, and it was confirmed that the product had desired strength and excellent workability.

On the other hand, No. in Table 2-2. Since Nos. 2 and 3 were out of the annealing temperature range of the heat treatment 2 and the heat treatment 1, respectively, No. No. 5 was out of the range of the cooling stop temperature T1 of the heat treatment 2 and the holding temperature in the second temperature range. Since steel plates outside the component value range were used for 7, 8 and 12, the tensile strength was 1300 MPa or more, the TS × T.El value was 18,000 MPa ·% or more, or the TS × λ value was 40,000 MPa ·. It was confirmed that it did not have any of% or more.

Claims (7)

By mass%
C: 0.10% or more and 0.20% or less,
Si: 0.01% or more and 2.50% or less,
Mn: 3.5% or more and 6.0% or less,
P: 0.100% or less,
S: 0.0500% or less,
Al: 0.01% or more and 0.50% or less and N: 0.010% or less,
Ingredient composition with the balance consisting of Fe and unavoidable impurities,
Area ratio to the entire steel sheet structure
Lower bainite is 30% or more and less than 80%,
Martensite including tempered martensite is 10% or more and less than 50%,
Residual austenite is 10% or more and 30% or less,
Polygonal ferrite is 10% or less (including 0%),
The average amount of Mn in the retained austenite is 7% or more in mass%,
It has a steel plate structure in which the average amount of Mn in the lower bainite, martensite and polygonal ferrite is 4% or less in mass%.
Tensile strength is 1300 MPa or more, tensile strength x total elongation is 18000 MPa ·% or more,
A high-strength steel plate having a tensile strength x hole expansion ratio of 40,000 MPa ·% or more. The composition of the components is further increased by mass%.
Cr: 0.005% or more and 1.000% or less,
V: 0.005% or more and 1.000% or less,
Ni: 0.005% or more and 1.000% or less,
The first aspect of claim 1, wherein Mo: 0.005% or more and 1.000% or less and Cu: 0.01% or more and 2.00% or less are selected from one or more. High-strength steel plate. The composition of the components is further increased by mass%.
The invention according to claim 1 or 2, wherein one or two selected from Ti: 0.005% or more and 0.100% or less and Nb: 0.005% or more and 0.100% or less are contained. High-strength steel plate. The composition of the components is further increased by mass%.
B: The high-strength steel plate according to any one of claims 1 to 3, which contains 0.0003% or more and 0.0050% or less. The composition of the components is further increased by mass%.
Any of claims 1 to 4, which comprises one or two selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less. The high-strength steel plate described in item 1. The composition of the components is further increased by mass%.
Sb: 0.200% or less,
Sn: 0.200% or less,
Ta: 0.100% or less,
W: 0.500% or less,
Mg: 0.0050% or less,
Zr: The high-strength steel plate according to any one of claims 1 to 5, which contains one kind or two or more kinds selected from 0.1000% or less. The method for manufacturing a high-strength steel sheet according to any one of claims 1 to 6.
A first heat treatment step of holding a steel piece having the above component composition for 600 s or more and 108,000 s or less in a temperature range of _{(Ac 1} transformation point + 20 ° C.) or more and Ac _{3 transformation point or less after hot rolling.}
The steel sheet after the first heat treatment step is cooled, and the steel sheet is cold-rolled to obtain a cold-rolled steel sheet.
Then, the cold-rolled steel sheet is annealed in the austenite single-phase region for 15 seconds or more and 1000 seconds or less, and then cooled to the first temperature region of (Ms point-100 ° C.) or more and less than the Ms point at an average cooling rate of 10 ° C./s or more. Then, the temperature is raised to a second temperature range of 250 ° C. or higher, (Bs point -20 ° C.) or 450 ° C., whichever is lower, and subsequently held in the second temperature range for 15 seconds or more and 1000 seconds or less. A method for producing a high-strength steel plate, which comprises a second heat treatment step.
Here, the Ms point is the martensitic transformation start temperature, and the Bs point is a value obtained by the following formula.
Bs (° C.) = 830-270 x [C%] -90 x [Mn%] -37 x [Ni%] -70 x [Cr%] -83 x [Mo%] However, [X%] is for steel sheets. Let it be the mass% of the component element X.

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