CN118139997A - Steel sheet, member, and method for producing same - Google Patents

Abstract

The present invention provides a steel plate and a member having high strength, high ductility, excellent stretch flange formability and good chemical conversion treatability, and a method for manufacturing the same. A steel plate, wherein the component composition is set to contain specific amounts of C, Si, Mn, P, S, sol.Al, and N in mass %, the total and the remainder structure of polygonal ferrite, upper bainite, retained γ, fresh martensite, tempered martensite, and lower bainite are set to a specific ratio, the number of fresh martensite particles with an equivalent circle diameter of less than 1.2 μm and the number of retained γ particles are set to a specific ratio, and the number of fresh martensite particles with an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more and the number of retained γ particles are set to a specific ratio.

Description

Steel plate, component and method for manufacturing the same

Technical Field

The present invention relates to a steel sheet which is suitable for use as a press-formed product having a complex shape after a press-forming process for use in automobiles, home appliances, etc. and has excellent chemical conversion treatability, a component using the steel sheet, and a method for producing the same.

Background Art

In the context of tightening _CO2 emission regulations around the world, there is a further demand for higher strength steel sheets for automobiles to reduce the weight of the vehicle body, and the application of higher strength steel sheets of 590MPa or higher is being promoted for vehicle body and seat parts, from the existing 440MPa-grade cold-rolled steel sheets. Generally, if the steel sheet is made higher strength, the stamping formability such as ductility and stretch flangeability will decrease, and cracks will be easily generated during stamping, and the freedom of shape will decrease. Therefore, it is limited to parts with simple shapes. Therefore, in order to apply high-strength steel sheets to parts with complex shapes, it is important to promote the high strength of steel sheets while maintaining or improving formability.

Under such circumstances, as a technology for improving the ductility of a steel sheet, TRIP steel in which retained austenite (retained γ) is dispersed in the microstructure of the steel sheet has been developed.

For example, Patent Document 1 discloses a manufacturing method based on an isothermal quenching treatment (a treatment in which a steel sheet containing C: 0.10 to 0.45%, Si: 0.5 to 1.8%, and Mn: 0.5 to 3.0% is cooled from a single-phase region annealing temperature or a dual-phase region annealing temperature to a bainite transformation temperature, is isothermally held, and thereby residual γ is formed by utilizing the bainite transformation during isothermal holding or cooling). The method discloses that a steel sheet containing C: 0.10 to 0.45%, Si: 0.5 to 1.8%, and Mn: 0.5 to 3.0% is subjected to an aging treatment in a temperature range of 350 to 500° C. for 1 to 30 minutes after annealing, thereby forming residual γ, and obtaining a steel sheet having a high ductility of TS: 80 kgf/mm ² or more and TS×EL: 2500 kgf/mm ² ·% or more.

Patent document 2 discloses that a steel plate containing C: 0.10-0.25%, Si: 1.0-2.0%, and Mn: 1.5-3.0% is cooled to 450-300°C at a rate of 10°C/second or more after annealing, and held for 180-600 seconds, thereby controlling the microstructure so as to obtain a steel plate having excellent ductility: El and stretch flangeability: λ, by having a retained γ volume ratio of 5% or more, bainitic ferrite of 60% or more, and polygonal ferrite of 20% or less.

In the above technology, in order to promote efficient carbon enrichment to the untransformed γ, the amount of Si contained in the steel plate is large. On the other hand, the thin steel plate used for stamping parts is later painted and assembled into automobiles, etc., so a chemical conversion treatment is performed for the purpose of giving the steel plate good paintability. During the chemical conversion treatment, in the presence of an oxide film on the surface of the steel plate, the crystalline particles attached by the chemical conversion treatment become uneven, which becomes a factor that deteriorates the paintability. Therefore, usually, pickling treatment is performed in a continuous annealing furnace used in the manufacture of thin steel plates as a pretreatment to improve the chemical conversion treatability, but in the composition of the steel plate, especially in the steel plate with a high Si content, the chemical conversion treatability is significantly deteriorated due to the Si-containing surface oxide layer that cannot be removed by pickling.

In order to solve such a problem, for example, Patent Document 3 discloses that after continuously immersing in a mixed acid solution containing an oxidizing first acid and a non-oxidizing second acid for pickling, it is continuously immersed in an acid solution containing a non-oxidizing third acid for re-pickling. Through such a process, a steel plate with a high Si content can have excellent chemical conversion treatability.

Prior art literature

Patent Literature

Patent Document 1: Japanese Patent Publication No. 6-35619

Patent Document 2: Japanese Patent No. 4411221

Patent Document 3: Japanese Patent No. 6041079

Summary of the invention

Problems to be solved by the invention

Simple shape parts can be formed by press forming only by processing to achieve uniform elongation, but for complex shape parts, local elongation is also important. However, although the conventional TRIP steel described in Patent Document 1 is excellent in E1, it has the problem of very low stretch flange formability.

In the technology described in Patent Document 2, bainitic ferrite is mainly used as the microstructure, and ferrite is suppressed to a small extent. Therefore, although the stretch flange formability is excellent, the ductility is not necessarily high. In addition, it is a technology related to a steel sheet with a low yield ratio, and it is difficult to apply it to vehicle body frame members and energy absorbing members. In addition, chemical conversion treatability is not considered in Patent Documents 1 and 2, and it is conceivable that the chemical conversion treatability deteriorates depending on the Si content or annealing conditions.

In the technology described in Patent Document 3, a steel sheet with high ductility and excellent stretch flangeability is obtained by utilizing the retention of the upper bainite transformation during the cooling process after annealing, the subsequent Q&P treatment, and the bainite transformation after reheating, but no improvement in the local elongation of the bending formability and bulging properties required for difficult-to-form parts has been confirmed. In addition, the carbon distribution from the martensite formed by the Q&P treatment to the untransformed γ is promoted, and a large amount of Si is contained, so the pickling technology described in Patent Document 3 is required. Therefore, it is necessary to build additional pickling equipment or the operating cost becomes high, so it is hoped that other technologies will be established.

As described above, the conventional technology has not been able to be said to be sufficient as a technology for a steel sheet that ensures high ductility and excellent stretch flange formability while having excellent chemical conversion treatability.

The present invention has been made to solve such problems, and its object is to provide a steel plate, a member and a method for manufacturing the same having a tensile strength of 590 MPa or more, achieving high ductility and excellent stretch flange formability and good chemical conversion treatability.

Here, the tensile strength of 590 MPa or more means that for a JIS No. 5 tensile test piece having a tensile direction in a direction perpendicular to the rolling direction, the tensile strength is 590 MPa or more in a tensile test in accordance with the provisions of JIS Z 2241 (2011) with the crosshead speed set to 10 mm/min.

In addition, high ductility means that for a JIS No. 5 tensile test piece having a tensile direction in a direction perpendicular to the rolling direction, the crosshead speed is set to 10 mm/min, and the tensile test is performed in accordance with the provisions of JIS Z 2241 (2011), the tensile strength (TS) × total elongation (T.El) is ≥ 22000 MPa·% or more.

In addition, excellent stretch flange formability means satisfying the following (A1) or (A2) by a hole expansion test in accordance with JFST 1001 (Japan Iron and Steel Federation Standard).

(A1) When the tensile strength is 590 MPa or more and less than 780 MPa, the hole expansion ratio λ is 60% or more.

(A2) When the tensile strength is 780 MPa or more, the hole expansion ratio λ is 35% or more.

In addition, good chemical conversion treatment property means that the steel plate is electrolytically pickled with sulfuric acid at a current density of 20 to 35 A/ ^dm2 for 2 seconds, degreased (treatment temperature 40°C, treatment time 120 seconds, spray degreasing), surface adjusted (pH 9.5, treatment temperature room temperature, treatment time 20 seconds), and then chemically converted using a zinc phosphate chemical conversion treatment solution (chemical conversion treatment solution temperature 35°C, treatment time 120 seconds), and there is no surface on which a chemical conversion film structure is not formed.

Methods used to solve problems

The present inventors have conducted intensive research on a method for achieving high ductility and excellent stretch flange formability even in a steel sheet composition with a low Si content, and have reached the following conclusions. Although not particularly limited, a low Si content refers to a case where the Si content is less than 1.60 mass %.

During austempering, carbon is distributed to the untransformed austenite until the free energy of the fcc phase and the bcc phase is equal to the _T0 composition through bainite transformation at about 400°C, and then the bainite transformation stops. Therefore, the coarse and thermally unstable untransformed austenite is transformed into a hard martensite structure or mechanically unstable retained γ during final cooling, and stretch flangeability is deteriorated. As such, it is generally difficult to achieve both ductility and stretch flangeability during austempering.

The inventors conducted an in-depth study on the heating process before annealing and found that by controlling the cold rolling conditions, steel composition and heating conditions, a soft ferrite structure and adjacent acicular austenite are formed through recrystallization during the heating process. The acicular austenite contributes to the carbon distribution in the structure formation during the cooling process and the formation of residual austenite.

In addition, in view of the fact that the lower the bainite transformation temperature, the more the _T0 composition expands to the high carbon region, a two-stage isothermal quenching treatment on both the high temperature side and the low temperature side has been studied. Thus, carbon distribution can be further promoted and uniform elongation can be improved. It is also found that by allowing the bainite transformation to proceed to the later stage without stopping the phase transformation, the amount and size of the coarse and hard fresh martensite generated in the final cooling are reduced. Moreover, since the needle-shaped austenite undergoes bainite transformation, it is possible to obtain a residual austenite with a high aspect ratio and high processing stability. It is thus found that stress concentration during stamping is suppressed, the formation of voids is suppressed, and thus the local elongation is improved. Even if such a needle-shaped austenite becomes a hard fresh martensite structure during the final cooling, the hole expansion will not be deteriorated.

The steel plate manufactured based on the above principles can improve both uniform elongation and local elongation even if the alloy design reduces the Si content. Therefore, there is no need for the expensive pickling treatment required to impart chemical conversion treatability by reducing the Si content, and a significant reduction in process costs can be achieved. It should be noted that the chemical conversion treatability recorded here refers to the characteristics that the adhesion amount and unevenness can meet the coating properties if it is a general pickling process. For example, the general pickling process refers to sulfuric acid pickling, etc., and the pickling method is not limited.

As described above, it was found that even a steel sheet with a low Si content exhibits excellent ductility and stretch flangeability. This is based on the following points.

(i) In the cold rolling process, by suppressing the development of shear texture in the first pass of cold rolling with a cold rolling ratio (reduction ratio) of 5% or more and less than 25%, a cold-rolled steel sheet in which a rolling preferential orientation (texture) and a Rotated Cube orientation are developed (a cold-rolled steel sheet having a total of structures of {111}<0-11> orientation, {111}<11-2> orientation, {211}<0-11> orientation and {100}<011> orientation of 35% or more and 75% or less in terms of area ratio relative to the total structure of the bcc phase) is manufactured.

(ii) In the annealing step, during the heating process before the uniform heat holding, the heating rate (average heating rate) from 500°C to Ac1 is set to less than 15°C/second, so that the structure of the cold-rolled steel sheet cold-rolled at a rolling ratio of more than 30% is fully recrystallized and the recrystallized texture is developed.

(iii) After that, the austenite (γ) that transforms at a temperature above Ac1 nucleates from the grain boundary of the recrystallized bcc phase or the residual carbide, but has a specific crystal orientation relationship relative to the surrounding bcc phase. Therefore, the interface matching is high, and the grain growth accompanied by the interface movement is delayed, but in order to approach the equilibrium state, a part of the interface moves preferentially to form needle-shaped austenite (needle-shaped γ). In order to make use of this needle-shaped γ, the annealing temperature is set to dual-phase annealing, and annealing is performed at an annealing temperature T that satisfies (T-Ac1)/(Ac3-Ac1)<1.0.

(iv) During cooling after isothermal holding in the annealing step, isothermal holding (first holding) is performed in a temperature range of 400 to 550°C, thereby transforming the needle-shaped γ into upper bainite with a small amount of precipitates and forming untransformed austenite (untransformed γ) with a high amount of dissolved C.

In addition, by a two-stage isothermal quenching treatment on the high temperature side and the low temperature side, carbon distribution is efficiently performed without stopping the bainite transformation, and needle-shaped austenite is formed in the structure before cooling. As a result, residual austenite with a high aspect ratio and high processing stability can be formed after the bainite transformation. As a result, it is possible to manufacture steel sheets that simultaneously achieve excellent uniform elongation and local elongation, and no cracks are generated even in stamping forming such as coexistence of bulging and stretch flange forming, and it can be applied to complex formed products.

Thus, by utilizing the acicular austenite formed by the heating process and the bainite transformation at two stages on the high temperature side and the low temperature side, excellent uniform elongation and local elongation can be obtained at the same time. As a result, even with component steels in which Si is suppressed to a small amount, a steel sheet having both high ductility and excellent stretch flange formability can be obtained, and a steel sheet having improved chemical conversion treatability can be obtained.

The present invention has been accomplished based on the above findings, and specifically provides the following solutions.

[1] A steel sheet having a composition comprising, by mass%, 0.06 to 0.24% C, 0.4% or more and less than 1.60% Si, 1.5 to 3.2% Mn, 0.02% or less P, 0.01% or less S, less than 1.0% sol.Al, less than 0.015% N, and satisfying the following formula (1), with the balance being Fe and unavoidable impurities,

The steel plate has a structure in which the area ratio of polygonal ferrite is 20% to 85%, the area ratio of upper bainite is 9% to 45%, the volume ratio of retained austenite is 3% to 15%, the area ratio of fresh martensite is 3% to 15%, the total area ratio of tempered martensite and lower bainite is 50% or less (including 0%), and the area ratio of residual structure is 5% or less,

The ratio of the total number of fresh martensite particles and retained austenite particles having an equivalent circle diameter of less than 1.2 μm to the total number of fresh martensite particles and retained austenite particles is 50% or more,

Furthermore, the ratio of the total number of fresh martensite particles and retained austenite particles having an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more to the number of fresh martensite particles and retained austenite particles having an equivalent circle diameter of 1.2 μm or more is 40% or more.

Si/Mn<0.50…Formula (1)

Here, in the formula (1), Si and Mn represent the Si content (mass %) and the Mn content (mass %), respectively.

[2] The steel sheet according to [1], further comprising, as the above-mentioned component composition, one or more selected from the group consisting of Nb: 0.2% or less, Ti: 0.2% or less, V: 0.2% or less, B: 0.01% or less, Cu: 0.2% or less, Ni: 0.2% or less, Cr: 0.4% or less, and Mo: 0.15% or less, in terms of mass %.

[3] The steel sheet according to [1] or [2], wherein the above-mentioned component composition further contains, in terms of mass%, one or more selected from the group consisting of Mg: 0.0050% or less, Ca: 0.0050% or less, Sn: 0.10% or less, Sb: 0.10% or less, and REM: 0.0050% or less.

[4] A member formed using the steel plate according to any one of [1] to [3] above.

[5] A method for manufacturing a steel plate, comprising:

a cold rolling step of hot-rolling and pickling the steel slab having the composition described in any one of [1] to [3] above, and then cold-rolling the obtained hot-rolled steel sheet to obtain a cold-rolled steel sheet; and

The annealing step is to anneal the cold-rolled steel sheet to obtain a steel sheet.

The cold rolling process includes the following steps:

The cumulative cold rolling ratio is set to 30 to 85%, and the first pass reduction ratio is set to 5% or more and less than 25%, thereby obtaining the above-mentioned cold-rolled steel sheet having a total of structures having a {111}<0-11> orientation, a {111}<11-2> orientation, a {211}<0-11> orientation, and a {100}<011> orientation of 35% or more and 75% or less in terms of area ratio relative to the total structure of the bcc phase,

The annealing process includes the following annealing treatment:

The cold rolled steel sheet is heated to an annealing temperature T of 840°C or less and satisfying 0.5≤(T-Ac1)/(Ac3-Ac1)<1.0, with an average heating rate of 0.5 to 15°C/second in a temperature range of 500°C or more and Ac1 or less.

After the heating, the steel sheet is held at the annealing temperature T in a furnace atmosphere having a dew point Td of -50°C or higher and -30°C or lower, thereby obtaining a steel sheet having a needle-like austenite structure with a number density of 5 or more per 1000 μm ² .

Next, the first cooling is performed by setting the average cooling rate in the temperature range of 750 to 550° C. to 6.0° C./sec or more, and cooling to a first cooling stop temperature Tc1 of 550° C. or less and 400° C. or more.

After the first cooling, a first holding is performed at the first cooling stop temperature Tc1 for more than 25 seconds.

After the first holding, a second cooling is performed to cool to a second cooling stop temperature Tc2 which is lower than the first cooling stop temperature Tc1 and lower than 450° C. and higher than 300° C.

A second holding period of 20 to 3000 seconds is performed at the second cooling stop temperature Tc2.

After the second holding, a third cooling step of cooling is performed.

[6] A method for manufacturing a component, comprising the step of subjecting the steel plate according to any one of [1] to [3] above to at least one of forming and joining to produce the component.

Effects of the Invention

According to the present invention, a steel sheet, a member and a method for producing the same are provided which have a tensile strength of 590 MPa or more, achieve high ductility and excellent stretch flange formability and good chemical conversion treatability.

The steel sheet of the present invention is suitable for use as press-formed products of complex shapes used in automobiles, home appliances, and the like after a press-forming process.

BRIEF DESCRIPTION OF THE DRAWINGS

1 is a diagram showing SEM photographs of a structure of acicular austenite (acicular γ) observed in a structure after final cooling (after third cooling in an annealing step) and a structure after being held at a temperature T and then water-cooled in the present invention.

FIG. 2 is a schematic diagram of needle-shaped austenite (acicular γ), and is a diagram for explaining the definition of the aspect ratio of the needle-shaped γ.

DETAILED DESCRIPTION

Hereinafter, the present invention will be described in detail. It should be noted that the present invention is not limited to the following embodiments.

The steel sheet of the present invention is a steel sheet as described below, which contains, by mass%, C: 0.06-0.24%, Si: 0.4% or more and less than 1.60%, Mn: 1.5-3.2%, P: 0.02% or less, S: 0.01% or less, sol.Al: less than 1.0%, N: less than 0.015%, and has a component composition satisfying the following formula (1), with the balance being Fe and inevitable impurities, and has an area ratio of polygonal ferrite of 20% or more and 85% or less, an area ratio of upper bainite of 9% or more and 45% or less, a volume ratio of retained austenite of 3% or more and 15% or less, and an area ratio of fresh martensite of 1.5% or more and 2.0% or less. The area ratio of the tempered martensite and the lower bainite is greater than 3% and less than 15%, the total area ratio of the tempered martensite and the lower bainite is less than 50% (including 0%), and the area ratio of the residual structure is less than 5%, the total number of fresh martensite particles and retained austenite particles with an equivalent circle diameter less than 1.2μm is more than 50% of the total number of fresh martensite particles and retained austenite particles, and the total number of fresh martensite particles and retained austenite particles with an aspect ratio of greater than 2.5 and an equivalent circle diameter of greater than 1.2μm is more than 40% of the total number of fresh martensite particles and retained austenite particles with an equivalent circle diameter of greater than 1.2μm.

Si/Mn<0.50…Formula (1)

Here, in the formula (1), Si and Mn represent the Si content (mass %) and the Mn content (mass %), respectively.

Hereinafter, the steel plate of the present invention will be described in the order of component composition and steel structure.

The steel sheet of the present invention contains the following components. In the following description, "%" as a unit of content of a component means "mass %".

C: 0.06～0.24%

C is contained from the viewpoint of making the hardenability of the steel plate, the strength of martensite and the volume ratio of residual γ within the desired range. When the C content is less than 0.06%, the strength and ductility of the steel plate cannot be fully ensured, so the C content is set to 0.06% or more. The C content is preferably 0.08% or more, more preferably 0.10% or more. When the C content exceeds 0.24%, the toughness of the weld deteriorates. In addition, when the C content exceeds 0.24%, the fresh martensite cannot be made to be the desired area ratio. Therefore, the C content is set to 0.24% or less. From the viewpoint of improving ductility and the toughness of the spot weld, the C content is preferably set to 0.21% or less. From the viewpoint of further improving the toughness of the spot weld, the C content is more preferably set to 0.20% or less.

Si: 0.4% or more and less than 1.60%

Si is contained from the viewpoint of achieving high strength of ferrite structure and stabilizing residual γ by suppressing the formation of carbides in martensite or bainite, thereby obtaining the effect of improving ductility. From these viewpoints, the Si content is set to 0.4% or more. From the viewpoint of improving ductility, the Si content is preferably set to 0.5% or more. The Si content is more preferably 0.6% or more. When the Si content is 1.60% or more, the chemical conversion treatability is significantly deteriorated. Therefore, the Si content is set to less than 1.60%. Preferably, the Si content is 1.30% or less, and more preferably 1.20% or less. It is further preferred that the Si content is less than 1.0%.

Mn: 1.5～3.2%

Mn ensures the specified hardenability, inhibits ferrite transformation, ensures the desired area ratio of tempered martensite and/or bainite, and thus ensures strength. In addition, Mn is enriched in γ during annealing in the dual phase region of ferrite/γ, thereby lowering the Ms point of the untransformed γ and stabilizing the residual γ, thereby improving ductility. In addition, Mn, like Si, inhibits the formation of carbides in bainite and improves ductility. In addition, Mn increases the volume ratio of residual γ and improves ductility. From these aspects, Mn is an important element in the present invention. In order to obtain these effects, the Mn content is set to 1.5% or more. From the viewpoint of improving hardenability, the Mn content is preferably set to 1.7% or more. The Mn content is more preferably set to 1.9% or more. On the other hand, when the Mn content exceeds 3.2%, the bainite transformation is significantly delayed, so it is difficult to ensure high ductility. In addition, when the Mn content exceeds 3.2%, it is difficult to suppress the formation of massive coarse γ, and the stretch flange formability is also deteriorated. Therefore, the Mn content is set to 3.2% or less. From the viewpoint of promoting bainite transformation and ensuring high ductility, the Mn content is preferably set to 3.0% or less, and more preferably set to 2.8% or less.

Si/Mn<0.50…Formula (1)

The surface oxide of the steel plate that significantly deteriorates the chemical conversion treatability is a Si-based oxide. Therefore, for the purpose of forming a Mn-containing oxide that is easily soluble in acid solution, Si/Mn is set to less than 0.50. That is, in the present invention, as formula (1), it is set to Si/Mn<0.50. Here, in formula (1), Si and Mn represent the Si content (mass %) and the Mn content (mass %), respectively. If within this range, it is possible to have chemical conversion treatability within the dew point range of above -50°C and below -30°C. Preferably, Si/Mn is below 0.40, and more preferably below 0.35.

P: 0.02% or less

P is an element that strengthens steel. When its content is high, the spot weldability deteriorates. Therefore, the P content is set to 0.02% or less. From the viewpoint of improving spot weldability, the P content is preferably set to 0.01% or less. It should be noted that P may not be contained, but from the viewpoint of manufacturing cost, the P content is preferably set to 0.001% or more.

S: 0.01% or less

S has the effect of improving the peeling property of oxide scale during hot rolling and the effect of suppressing nitridation during annealing, but S is an element that deteriorates the local elongation in addition to the spot weldability. In order to suppress these deteriorations, the S content is set to 0.01% or less. In the present invention, the contents of C, Si, and Mn are high, so the spot weldability is easily deteriorated. From the viewpoint of improving the spot weldability, the S content is preferably set to 0.0020% or less, and more preferably set to less than 0.0010%. It should be noted that S may not be contained, but from the viewpoint of manufacturing cost, the S content is preferably set to 0.0001% or more.

Sol.Al: less than 1.0%

Al is contained for the purpose of deoxidation or for the purpose of stabilizing the residual γ as a substitute for Si. The lower limit of sol.Al is not particularly specified, but is preferably set to 0.01% or more for stable deoxidation. On the other hand, when the sol.Al content is 1.0% or more, the strength of the raw material is extremely reduced, and the chemical conversion treatability is also deteriorated. In addition, a large amount of aluminum oxides are generated during steelmaking, thereby significantly deteriorating the bendability. Therefore, the sol.Al content is set to less than 1.0%. In order to obtain high strength, the sol.Al content is preferably set to less than 0.50%, and more preferably set to 0.10% or less.

N: less than 0.015%

N is an element that forms nitrides such as BN, AlN, and TiN in steel, and is an element that reduces the hot ductility of steel and reduces the surface quality. In addition, in steel containing B, there is a disadvantage that the effect of B disappears due to the formation of BN. When the N content is 0.015% or more, the surface quality deteriorates significantly. Therefore, the N content is set to less than 0.015%. It should be noted that N may not be contained, but from the perspective of manufacturing cost, the N content is preferably set to 0.0001% or more.

The balance other than the above is Fe and inevitable impurities. The steel sheet of the present invention preferably has a composition comprising the above-mentioned basic components and the balance being Fe and inevitable impurities.

The steel sheet of the present invention may contain one or two selected from the following (A) and (B) as the following optional elements in addition to the above components, instead of part of the above Fe and inevitable impurities.

(A) one or more selected from the group consisting of, in mass%, Nb: 0.2% or less, Ti: 0.2% or less, V: 0.2% or less, B: 0.01% or less, Cu: 0.2% or less, Ni: 0.2% or less, Cr: 0.4% or less, and Mo: 0.15% or less

(B) one or more selected from the group consisting of Mg: 0.0050% or less, Ca: 0.0050% or less, Sn: 0.10% or less, Sb: 0.10% or less, REM: 0.0050% or less, Nb: 0.2% or less, in terms of mass %

It is preferred to add Nb from the viewpoint of improving the defect resistance of the spot weld by refining the microstructure. In addition, Nb can be contained from the effect of refining the steel structure and increasing the strength, promoting the bainite transformation by graining, improving the bendability, and improving the delayed fracture resistance. In order to obtain these effects, the lower limit is not particularly specified, and the Nb content is preferably 0.002% or more. The Nb content is more preferably 0.004% or more, and further preferably 0.010% or more. However, when a large amount of Nb is contained, precipitation strengthening becomes too strong and ductility decreases. In addition, it leads to an increase in rolling load and deterioration of castability. Therefore, when Nb is contained, the Nb content is set to 0.2% or less. The Nb content is preferably 0.1% or less, more preferably 0.05% or less, and further preferably 0.03% or less.

Ti: 0.2% or less

It is preferred to add Ti from the viewpoint of improving the defect resistance of the spot weld by refining the microstructure. In addition, Ti has the effect of fixing the N in the steel in the form of TiN to produce the effect of improving the hot ductility and the effect of improving the hardenability of B. In order to obtain these effects, the lower limit is not particularly specified, and the Ti content is preferably set to more than 0.002%. From the viewpoint of fully fixing N, the Ti content is further preferably more than 0.008%. More preferably more than 0.010%. On the other hand, when the Ti content exceeds 0.2%, the increase in rolling load and the increase in precipitation strengthening amount cause a decrease in ductility. Therefore, when containing Ti, the Ti content is set to less than 0.2%. The Ti content is preferably less than 0.1%, and more preferably less than 0.05%. In order to ensure high ductility, the Ti content is further preferably set to less than 0.03%.

V: 0.2% or less

V can be contained from the effect of improving the hardenability of steel, the effect of suppressing the formation of carbides in martensite and upper bainite/lower bainite, the effect of refining the structure, and the effect of precipitating carbides to improve the delayed fracture resistance. In order to obtain these effects, the lower limit is not particularly specified, and the V content is preferably set to 0.003% or more. The V content is more preferably 0.005% or more, and further preferably 0.010% or more. On the other hand, when a large amount of V is contained, the castability is significantly deteriorated. Therefore, when V is contained, the V content is set to 0.2% or less. The V content is preferably 0.1% or less. The V content is more preferably 0.05% or less, and further preferably 0.03% or less.

B: 0.01% or less

B has the advantage of easily generating tempered martensite and/or bainite of a specified area ratio. In addition, the delayed fracture resistance is improved by the residual solid solution B. In order to obtain such an effect of B, it is preferred to set the B content to 0.0002% or more. In addition, the B content is more preferably 0.0005% or more. The B content is further preferably 0.0010% or more. On the other hand, when the B content exceeds 0.01%, not only is its effect saturated, but it also leads to a significant decrease in hot ductility and the generation of surface defects. Therefore, when B is contained, the B content is set to 0.01% or less. The B content is preferably 0.0050% or less. The B content is more preferably 0.0030% or less.

Cu: 0.2% or less

Cu improves the corrosion resistance of automobiles in the use environment. In addition, the corrosion products of Cu cover the surface of the steel plate and have the effect of inhibiting the intrusion of hydrogen into the steel plate. Cu is an element mixed in when scrap steel is utilized as a raw material. By allowing the mixing of Cu, recycled materials can be utilized as raw materials, which can reduce manufacturing costs. Therefore, there is no special lower limit, and from the viewpoint of improving delayed fracture resistance, Cu is preferably contained at least 0.05%. The Cu content is more preferably at least 0.10%. On the other hand, when the Cu content is too much, surface defects will occur. Therefore, when Cu is contained, the Cu content is set to less than 0.2%.

Ni: 0.2% or less

Ni is also an element that has the effect of improving corrosion resistance like Cu. In addition, Ni has the effect of suppressing the generation of surface defects that are easy to occur when Cu is contained. In order to obtain such an effect, the lower limit is not particularly specified, and the Ni content is preferably set to 0.01% or more. The Ni content is more preferably 0.04% or more, and further preferably 0.06% or more. On the other hand, when the Ni content is too much, the oxide scale generation in the heating furnace becomes uneven, which becomes the cause of surface defects. In addition, it also leads to increased costs. Therefore, when Ni is contained, the Ni content is set to 0.2% or less.

Cr: 0.4% or less

Cr can be contained in order to improve the hardenability of steel and inhibit the formation of carbides in martensite and upper bainite/lower bainite. In order to obtain such an effect, the lower limit is not particularly specified, and the Cr content is preferably set to 0.01% or more. The Cr content is more preferably 0.03% or more, and further preferably 0.06% or more. On the other hand, when excessive Cr is contained, pitting resistance deteriorates. Therefore, when Cr is contained, the Cr content is set to 0.4% or less.

Mo: 0.15% or less

Mo can be contained from the effect of improving the hardenability of steel and the effect of suppressing the formation of carbides in martensite and upper bainite/lower bainite. In order to obtain such effects, the Mo content is preferably set to 0.01% or more. The Mo content is more preferably 0.03% or more, and further preferably 0.06% or more. On the other hand, Mo significantly deteriorates the chemical conversion treatability of the cold-rolled steel sheet, so when Mo is contained, the Mo content is set to 0.15% or less.

Mg: 0.0050% or less

Mg fixes O in the form of MgO, which contributes to the improvement of delayed fracture resistance. Therefore, the Mg content is preferably set to 0.0002% or more. The Mg content is more preferably 0.0004% or more, and further preferably 0.0006% or more. On the other hand, when a large amount of Mg is added, the surface quality and bendability deteriorate. Therefore, when Mg is contained, the Mg content is set to 0.0050% or less. The Mg content is preferably 0.0025% or less, and more preferably 0.0010% or less.

Ca: 0.0050% or less

Ca fixes S in the form of CaS, which contributes to the improvement of bendability and delayed fracture resistance. Therefore, the Ca content is preferably set to 0.0002% or more. The Ca content is more preferably 0.0005% or more, and further preferably 0.0010% or more. On the other hand, when Ca is added in large amounts, the surface quality and bendability deteriorate. Therefore, when Ca is contained, the Ca content is set to 0.0050% or less. The Ca content is preferably 0.0035% or less, and more preferably 0.0020% or less.

Sn: 0.10% or less

Sn inhibits oxidation and nitridation of the surface layer of the steel plate, and inhibits the reduction of the content of C and B in the surface layer caused by this. In addition, by inhibiting the above-mentioned reduction in the content of C and B, the ferrite formation in the surface layer of the steel plate is inhibited, high strength is achieved, and fatigue resistance is improved. From this point of view, the Sn content is preferably set to 0.002% or more. The Sn content is more preferably 0.004% or more, and further preferably 0.006% or more. The Sn content is more preferably 0.008% or more, and further preferably 0.01% or more.

On the other hand, when the Sn content exceeds 0.10%, the castability deteriorates. In addition, Sn segregates at the original γ grain boundary, and the delayed fracture resistance deteriorates. Therefore, when Sn is contained, the Sn content is set to 0.10% or less. The Sn content is preferably 0.04% or less, and more preferably 0.03% or less.

Sb: 0.10% or less

Sb inhibits oxidation and nitridation of the surface of the steel plate, and inhibits the resulting decrease in the content of C and B in the surface. In addition, by inhibiting the above-mentioned decrease in the content of C and B, the formation of ferrite in the surface of the steel plate is inhibited, high strength is achieved, and fatigue resistance is improved. From this point of view, the Sb content is preferably set to 0.002% or more. The Sb content is more preferably 0.004% or more, and further preferably 0.006% or more. On the other hand, when the Sb content exceeds 0.10%, the castability deteriorates, and in addition, it segregates at the original γ grain boundary, and the delayed fracture resistance deteriorates. Therefore, when Sb is contained, the Sb content is set to 0.10% or less. The Sb content is preferably 0.04% or less, and more preferably 0.03% or less.

REM: 0.0050% or less

REM is an element that suppresses the adverse effects of sulfides on stretch flange formability by spheroidizing the shape of sulfides, thereby improving stretch flange formability. In order to obtain these effects, the REM content is preferably set to 0.0005% or more. The REM content is more preferably 0.0010% or more, and even more preferably 0.0020% or more.

On the other hand, when the REM content exceeds 0.0050%, the effect of improving the stretch flange formability is saturated. Therefore, when REM is contained, the REM content is made 0.0050% or less.

It should be noted that the REM described in the present invention refers to scandium (Sc) with an atomic number of 21, yttrium (Y) with an atomic number of 39, and lanthanide elements ranging from lanthanum (La) with an atomic number of 57 to lutetium (Lu) with an atomic number of 71. The REM concentration in the present invention refers to the total content of one or more elements selected from the above REM.

When the optional components are contained in an amount less than the lower limit, the optional elements contained in an amount less than the lower limit do not impair the effects of the present invention. Therefore, when the optional elements are contained in an amount less than the lower limit, the optional elements are contained as unavoidable impurities.

Next, the steel structure of the steel plate of the present invention will be described.

Area ratio of polygonal ferrite: 20% or more and 85% or less

In order to ensure high ductility, the area ratio of polygonal ferrite is set to 20% or more. The area ratio of polygonal ferrite is preferably 25% or more, more preferably 30% or more. The area ratio of polygonal ferrite is preferably 35% or more, more preferably 40% or more.

On the other hand, in order to obtain a predetermined strength, the area ratio of polygonal ferrite is set to 85% or less. More preferably, the area ratio of polygonal ferrite is 82% or less. The area ratio of polygonal ferrite is preferably 80% or less, and more preferably 78% or less.

Area ratio of upper bainite: 9% or more and 45% or less

Upper bainite is bainite with less carbide precipitation, and in order to distribute C in the surrounding untransformed γ, the upper bainite can be used to form residual γ with high processing stability. In addition, the upper bainite has a hardness between ferrite and martensite, and the local elongation is improved by forming these intermediate hardness structures. Therefore, in the case of a strength level with a tensile strength (TS) of 590MPa or more, more than 9% of upper bainite is required. Therefore, the upper bainite is set to be 9% or more in terms of area ratio. Preferably, the area ratio of the upper bainite is 12.0% or more, and more preferably 15.0% or more. On the other hand, in the present invention in which the temperature is maintained at the dual phase region, in view of the formation of a large amount of ferrite structure, the area ratio of the upper bainite is set to 45% or less for the purpose of suppressing the reduction in strength. The upper bainite is preferably 38% or less, and more preferably 30% or less.

Volume ratio of retained austenite (retained γ): 3% or more and 15% or less

In order to ensure high ductility, the volume fraction of residual γ is set to be 3% or more relative to the entire steel structure. The volume fraction of residual γ (residual γ amount) is preferably 3.0% or more, more preferably 5% or more, and further preferably 7% or more. The residual γ amount includes both the residual γ generated adjacent to the upper bainite and the residual γ generated adjacent to the martensite and lower bainite. When the residual γ amount increases excessively, it leads to a decrease in strength and a significant decrease in stretch flange formability. Therefore, the volume fraction of residual γ is set to 15% or less. The volume fraction of residual γ is preferably 13% or less. In addition, "volume fraction" can be regarded as "area fraction".

Area ratio of fresh martensite: 3% or more and 15% or less

Fresh martensite is a structure that reduces local elongation, and by forming fresh martensite within a range that does not deteriorate bendability and hole expandability, strength can be improved. From this viewpoint, the area ratio of fresh martensite is set to have a lower limit of 3% and an upper limit of 15%.

Total area ratio of tempered martensite and lower bainite: 50% or less (including 0%)

In the present invention, tempered martensite and lower bainite are formed by maintaining the low temperature side of two-stage austenite tempering.

Compared with the upper bainite with less precipitation of carbides, the tempered martensite and lower bainite in which carbides are precipitated in the structure suppress the carbon distribution to the untransformed γ. However, the tempered martensite and lower bainite lead to the carbon enrichment to the untransformed γ caused by the expansion of the _T0 composition at low temperature, or further reduce the amount of fresh martensite during the final cooling, so it is necessary to control these structures to obtain the retained γ with high processing stability.

According to the steel component composition and the austenite tempering temperature, the low temperature side of the above two stages of austenite tempering is maintained below the Ms point, so a part of the untransformed γ undergoes martensitic transformation and is tempered by the subsequent maintenance. When the total area ratio of tempered martensite and lower bainite exceeds 50%, carbide precipitation is promoted, the necessary residual γ amount cannot be obtained, and the strength becomes too high, so the desired ductility cannot be obtained. Therefore, in the present invention, the total area ratio of tempered martensite and lower bainite is set to 50% or less. Preferably, the total area ratio is 40% or less, and more preferably 35% or less. The total area ratio is more preferably 30% or less, and even more preferably 25% or less.

On the other hand, if the area ratios of polygonal ferrite, upper bainite, retained γ, and fresh martensite can be controlled within a desired range, the total area ratio of tempered martensite and lower bainite can be 0%.

Area ratio of residual tissue: less than 5%

The residual structure is a structure other than polygonal ferrite, upper bainite, residual austenite, fresh martensite, tempered martensite, and lower bainite, and pearlite can be cited as an example. If a pearlite structure is formed, efficient carbon distribution is hindered and the formation of residual γ is inhibited, so the ductility is reduced. In addition, in the annealing process of implementing a two-stage isothermal quenching treatment, pearlite transformation occurs from austenite (γ) in a part of the structure according to the steel component composition and the austenite tempering temperature. In the present invention, if the area ratio of the residual structure is less than 5%, the influence on the material can be ignored, so the area ratio of the residual structure is set to an upper limit of 5%. The area ratio of the residual structure is preferably less than 5.0%. The area ratio of the residual structure can also be 0%.

The ratio of the total number of fresh martensite particles and retained γ particles with an equivalent circle diameter of less than 1.2 μm to the total number of fresh martensite particles and retained austenite particles: 50% or more

Fresh martensite particles and residual γ particles with an equivalent circle diameter of less than 1.2 μm are structures that are not likely to cause stress concentration during local deformation and do not contribute to the formation of voids, thereby not deteriorating local ductility and hole expandability.

If the total number of fresh martensite particles and retained γ particles having an equivalent circle diameter of less than 1.2 μm is 50% or more relative to the total number of fresh martensite particles and retained γ particles, excellent local elongation and hole expandability can be obtained in the present invention.

Therefore, in the present invention, the ratio of the total number of fresh martensite particles and retained γ particles having an equivalent circle diameter of less than 1.2 μm to the total number of fresh martensite particles and retained austenite particles is set to 50% or more. That is, the following formula (A) is satisfied.

100×(total number of fresh martensite particles and residual γ particles with equivalent circle diameter less than 1.2 μm)/(total number of fresh martensite particles and residual γ particles)≥50(%)…Formula (A)

The ratio of the left side defined in the above formula (A) is preferably 55% or more.

It should be noted that, in order to obtain the above-mentioned structure, upper bainite, tempered martensite and lower bainite can be formed in the structure through a two-stage austenite tempering process.

Ratio of the total number of fresh martensite particles and retained γ particles with an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more to the number of fresh martensite particles and retained γ particles with an equivalent circle diameter of 1.2 μm or more: 40% or more

In fresh martensite particles and/or retained γ particles with an equivalent circle diameter of 1.2 μm or more, by increasing the aspect ratio of the fresh martensite particles and/or retained γ particles, stress concentration during local deformation can be reduced, void formation can be suppressed, and local ductility and hole expandability can be improved.

For such fresh martensite particles and/or residual γ particles, the area ratio can be increased by causing the needle-shaped γ surrounded by the soft ferrite structure formed during the heating process to undergo bainite transformation during the subsequent cooling process. In the present invention, the desired formability can be achieved by making the particles with an equivalent circle diameter of 1.2 μm or more and an aspect ratio of 2.5 or more relative to the total number of fresh martensite particles and residual γ particles with an equivalent circle diameter of 1.2 μm or more 40%.

Therefore, in the present invention, the ratio of the total number of fresh martensite particles and retained austenite particles having an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more to the number of fresh martensite particles and retained austenite particles having an equivalent circle diameter of 1.2 μm or more is set to 40% or more. That is, together with the above formula (A), the following formula (B) is further satisfied. It is preferred that the ratio on the left side specified in the following formula (B) is 45% or more.

100×(the total number of fresh martensitic particles and residual γ particles with an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more)/(the number of fresh martensitic particles and residual γ particles with an equivalent circle diameter of 1.2 μm or more)≥40…Formula (B)

The structure of the obtained steel plate was measured by the following method.

Determination of Area Ratio of Steel Structure

The observation sample was cut from the steel plate in a manner that the cross section perpendicular to the steel plate surface and parallel to the rolling direction was used as the observation surface, and the plate thickness cross section was corroded with a 1 volume % nitric acid ethanol solution. The scanning electron microscope (SEM) was magnified 2000 times, and the tissue photograph was taken in the area of 3000 μm ² or more at the plate thickness t/4. Then, the following items (i) to (iv) were measured respectively. It should be noted that t represents the plate thickness and w represents the plate width.

(i) Polygonal ferrite and upper bainite

Polygonal ferrite (recrystallized F) and upper bainite (UB) both appear gray in the SEM photograph, but can be identified by their shapes. An example of an SEM photograph is shown in FIG1 together with an SEM photograph of the structure after water cooling after being held at temperature T. The area represented by the dotted line in FIG1(a) is a needle-shaped γ structure formed by a treatment of isothermal holding at an annealing temperature T within the scope of the present invention in the annealing process, in which upper bainite (UB) is generated, and residual γ or fresh martensite (M) with a high aspect ratio is formed around it. The same structure is also observed in the blocky γ structure formed by isothermal holding at an annealing temperature T. The area ratios of polygonal ferrite and upper bainite are measured by the point algorithm according to ASTM E562-11 (2014). The area ratio of polygonal ferrite and the area ratio of upper bainite are the average values of the measured values at 5 locations, respectively.

(ii) Fresh martensite and residual γ

Both fresh martensite and residual γ appear white in the SEM photograph and cannot be distinguished. Therefore, residual γ is measured separately by the method described later. In addition, the total area ratio of fresh martensite and residual γ is measured by the point algorithm according to ASTME562-11 (2014) from the SEM photograph, and the area ratio of residual γ measured by the method described later is subtracted from the total area ratio to determine the area ratio of fresh martensite. The total area ratio of fresh martensite and residual γ is measured by the point algorithm, and the value obtained by subtracting the volume ratio of residual γ measured by the method described later from the average value of the measured values at 5 places is set as the area ratio of fresh martensite.

(iii) Tempered martensite and/or lower bainite

Tempered martensite and lower bainite are structures containing carbides observed as white fine structures in SEM photos. The two can be distinguished in a more microscopic observation, but it is difficult to distinguish in SEM photos. Therefore, in the present invention, tempered martensite and lower bainite are defined as the same structure, and the total area ratio of tempered martensite and lower bainite is measured by point calculation according to ASTM E562-11 (2014). The average value of the measured values at 5 locations is used as the total area ratio of tempered martensite and lower bainite.

(iv) Surplus organization

The area ratio of the remaining structure is defined as the area ratio of the remaining structure by subtracting the area ratios of polygonal ferrite, upper bainite, fresh martensite, retained γ, tempered martensite and lower bainite measured by the above method from 100%.

Determination of volume fraction of retained γ

After the steel plate is ground to a position of 1/4 of the plate thickness, it is further ground by chemical grinding by 0.1 mm. For the surface thus obtained, the integrated reflection intensity of the (200) plane, (220) plane, (311) plane of FCC iron (γ) and the (200) plane, (211) plane, (220) plane of BCC iron (ferrite) is measured by using Mo Kα rays using an X-ray diffraction device, and the volume ratio of residual γ obtained from the intensity ratio of the integrated reflection intensity from each surface of FCC iron (γ) to the integrated reflection intensity from each surface of BCC iron (ferrite) is measured. In the present invention, the volume ratio of the residual γ can be used as the area ratio of the residual γ.

Equivalent circle diameter and aspect ratio of fresh martensitic particles and/or residual γ particles

An observation sample was cut from the steel plate with the cross section parallel to the rolling direction as the observation surface, and the structural corrosion of the plate thickness cross section was visualized using Lepera etching solution. A laser microscope (LM) was used to magnify the area to 1000 times at a thickness of t/4 and above 10000 ^μm2 to take a structural photograph.

Lepera etching is a color etching method that represents fresh martensite and/or residual γ with a white contrast, extracts fresh martensite particles and/or residual γ particles, and performs image analysis to determine the equivalent circular diameter and aspect ratio of the fresh martensite particles and/or residual γ particles.

Among all the obtained particles, the number of particles having an equivalent circle diameter of less than 1.2 μm was measured, and the ratio of the number of particles to the number of all the particles was calculated.

In addition, among all the particles obtained, particles with an equivalent circle diameter of 1.2 μm or more are taken as the object, and the number of particles with an aspect ratio of 2.5 or more is measured, and the ratio of particles with an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more to all particles with an equivalent circle diameter of 1.2 μm or more is calculated.

Next, an embodiment of the method for manufacturing a steel plate of the present invention is described in detail. It should be noted that, unless otherwise specified, the temperatures when heating or cooling a steel slab (steel raw material), steel plate, etc. shown below refer to the surface temperatures of the steel slab (steel raw material), steel plate, etc.

The method for manufacturing a steel sheet of the present invention comprises: a cold rolling process of performing hot rolling and pickling on a steel billet having the above-mentioned composition, and then performing cold rolling on the obtained hot-rolled steel sheet to obtain a cold-rolled steel sheet; and an annealing process of performing annealing on the cold-rolled steel sheet to obtain a steel sheet, wherein the cold rolling process comprises the following cold rolling process: setting the cumulative cold rolling ratio to 30 to 85%, setting the first pass reduction ratio to 5% or more and less than 25%, thereby obtaining a total phase of a structure having a {111}<0-11> orientation, a {111}<11-2> orientation, a {211}<0-11> orientation, and a {100}<011> orientation. For a cold-rolled steel sheet in which the total structure of the bcc phase is 35% or more and 75% or less in terms of area ratio, the annealing step includes the following annealing treatment: for the cold-rolled steel sheet, the average heating rate in the temperature range of 500°C or more and Ac1 or less is set to 0.5 to 15°C/second, and the steel sheet is heated to an annealing temperature T of 840°C or less and satisfying 0.5≤(T-Ac1)/(Ac3-Ac1)<1.0, and after the heating, the steel sheet is held at the annealing temperature T in a furnace atmosphere having a dew point Td of -50°C or more and -30°C or less, thereby obtaining a needle-shaped austenite structure with a number density of 5 pieces/1000μm ² or more, then, a first cooling is performed in which the average cooling rate in the temperature range of 750 to 550°C is set to 6.0°C/second or more, and the steel plate is cooled to a first cooling stop temperature Tc1 of 550°C or less and 400°C or more. After the first cooling, a first holding is performed in which the steel plate is kept at the first cooling stop temperature Tc1 for 25 seconds or more. After the first holding, a second cooling is performed in which the steel plate is cooled to a second cooling stop temperature Tc2 of 450°C or less and 300°C or more and the first cooling stop temperature Tc1 or less is carried out. A second holding is performed in which the steel plate is kept at the second cooling stop temperature Tc2 for 20 to 3000 seconds. After the second holding, a third cooling is performed in which the steel plate is cooled.

Hereinafter, each step will be described.

Hot rolling process

In the present invention, as hot rolling in the hot rolling process, there are methods of rolling after reheating the steel billet having the above-mentioned composition, methods of directly rolling the steel billet after continuous casting without heating, methods of rolling after short-time heating treatment of the steel billet after continuous casting, etc. Hot rolling can be carried out according to conventional methods. For example, the steel billet heating temperature can be set to 1100°C or more and 1300°C or less, the soaking time can be set to 20 to 30 minutes, the finishing temperature can be set to Ar3 transformation point (°C) or more and Ar3 transformation point (°C) + 200°C or less, and the coiling temperature can be set to 400 to 720°C. From the viewpoint of suppressing the variation of the plate thickness and stably ensuring high strength, the coiling temperature is preferably set to 430 to 620°C.

The melting method for producing the above-mentioned steel slab (steel raw material) is not particularly limited, and a known melting method such as a converter or an electric furnace can be used. In addition, secondary refining can also be performed using a vacuum degassing furnace.

Pickling process

The pickling treatment step is a step of performing pickling treatment on the hot-rolled steel sheet after the hot rolling step. The pickling treatment conditions are not particularly limited, and the pickling treatment conditions in a known production method may be adopted.

Cold rolling process

Cumulative cold rolling rate: 30～85%

When the reduction rate (cumulative cold rolling rate) in the cold rolling treatment is less than 30%, recrystallization cannot be sufficiently promoted, and the formation of the needle-shaped γ mentioned in the present invention cannot be fully performed. In addition, the structure after the annealing process becomes uneven. Therefore, the reduction rate (cumulative cold rolling rate) of cold rolling needs to be carried out in a range of more than 30%, preferably more than 40%, and more preferably more than 50%. On the other hand, from the perspective of cold rolling load or further from the perspective of material, the reduction rate (cumulative cold rolling rate) is less than 85%.

In the cold rolling process, the number of passes is not particularly limited, and can be set to 5 passes, for example. The cumulative cold rolling rate (plate thickness reduction rate) is (1-(plate thickness after cold rolling (after the final pass)/plate thickness before cold rolling)×100.

The first pass reduction rate: 5% or more and less than 25%

The first pass reduction ratio is set to 5% or more from the viewpoint of operability. On the other hand, when the first pass reduction ratio is 25% or more, since the sheet temperature during the first pass cold rolling is low, the cold rolled material is given a shear component strain, the desired texture is not developed, and the needle-shaped γ is not formed. Therefore, the first pass reduction ratio is set to 5% or more and less than 25%.

The first-pass rolling reduction rate (plate thickness reduction rate) refers to (1-(plate thickness after the first-pass cold rolling)/(plate thickness before cold rolling))×100.

The rolling temperature (plate temperature) of the first pass is preferably 20°C or more and 40°C or less. It should be noted that the rolling temperature of the first pass is obtained by measuring the portion of the steel plate surface to which the lubricating oil is not attached after one pass using a radiation thermometer. When the rolling temperature of the first pass is lower than 20°C or when the rolling temperature of the first pass exceeds 40°C, sometimes the desired texture is not developed and needle-shaped γ is not formed. Therefore, the rolling temperature of the first pass is preferably 20°C or more and 40°C or less.

The structure of the cold-rolled steel sheet after cold rolling: the total of the structures having {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation and {100} <011> orientation is 35% or more and 75% or less in terms of area ratio relative to the total structure of the bcc phase

The needle-shaped γ has a specific crystal orientation relationship with the ferrite around its nucleation site (Near Kurdjumov-Sachs relationship).

In the structure of the cold-rolled steel sheet after cold rolling, by making the total of structures having prescribed orientations such as {111}<0-11> orientation, {111}<11-2> orientation, {211}<0-11> orientation and {100}<011> orientation a certain amount or more in terms of area ratio relative to the total structure of the bcc phase, the reverse transformation γ having the above prescribed orientation is easily formed between the surrounding ferrite grains, resulting in the formation of a large amount of needle-shaped γ. In order to form the desired amount of needle-shaped γ, it is necessary to make the total of structures having {111}<0-11> orientation, {111}<11-2> orientation, {211}<0-11> orientation and {100}<011> orientation 35% or more in terms of area ratio relative to the total structure of the bcc phase. Preferably, it is 40% or more. On the other hand, when the total of the structures having the {111}<0-11> orientation, the {111}<11-2> orientation, the {211}<0-11> orientation, and the {100}<011> orientation exceeds 75% by area ratio relative to the total structure of the bcc phase, material anisotropy of the steel sheet occurs. Therefore, the total of the structures having the above-mentioned prescribed orientations is set to 75% or less by area ratio relative to the total structure of the bcc phase. It is preferably 70% or less, and more preferably 65% or less.

In the present invention, by cold rolling the hot-rolled steel sheet having the above-mentioned composition at a cold rolling rate of 30 to 85%, and setting the reduction rate of the first pass to be greater than 5% and less than 25%, the ratio of the total area ratio of the structure having the above-mentioned specified orientation to the area ratio of the total structure of the bcc phase can be adjusted to a desired range.

Determination method of texture of cold rolled structure

A measurement sample having a cross section parallel to the rolling direction as a measurement surface is cut from the cold-rolled steel sheet after the cold rolling process, and the measurement surface is mechanically polished or electrolytically polished, and then the area of 80000 μm ² or more is measured by the SEM-EBSD method (measurement conditions: WD: 20 mm, acceleration voltage: 20 kV). The area ratio of the bcc phase with the {ND plane} <RD direction> orientation of the rolling {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100} <011> orientation is quantified, and the ratio to the area ratio of the bcc phase of all orientations is calculated, thereby evaluating the texture of the cold-rolled steel sheet.

Annealing process

The annealing process of the present invention is an annealing treatment as described below: for the cold-rolled steel sheet after the above-mentioned cold rolling process, the average heating rate (HR1) in the temperature range of 500°C to Ac1 is set to 0.5 to 15°C/second, and heating is performed until the annealing temperature T is 840°C or less and satisfies 0.5≤(T-Ac1)/(Ac3-Ac1)<1.0, and after the heating, the annealing temperature T is maintained at the above-mentioned annealing temperature T in a furnace atmosphere with a dew point Td of -50°C to -30°C, thereby obtaining a needle-shaped austenite structure with a number density of 5 pieces/1000μm ² or more, then, a first cooling is performed in which the average cooling rate in the temperature range of 750 to 550°C is set to 6.0°C/second or more, and the steel plate is cooled to a first cooling stop temperature Tc1 of 550°C or less and 400°C or more. After the first cooling, a first holding is performed in which the first cooling stop temperature Tc1 is maintained for 25 seconds or more. After the first holding, a second cooling is performed in which the steel plate is cooled to a second cooling stop temperature Tc2 of less than the first cooling stop temperature Tc1 and less than 450°C and more than 300°C. A second holding is performed in which the steel plate is maintained at the second cooling stop temperature Tc2 for 20 to 3000 seconds. After the second holding, a third cooling is performed in which cooling is performed.

Average heating rate in the temperature range of 500°C or higher and Ac1 or lower: 0.5 to 15°C/sec or lower

In the present invention, the cold-rolled sheet having the structure after the cold rolling process is heated at an appropriate heating rate to fully promote recrystallization, and then the needle-shaped γ is formed by heating to temperature T or holding at temperature T. Therefore, in the temperature range of Ac1 or below where γ transformation does not occur at 500°C or above, the average heating rate is set to 15°C/second or less. The average heating rate is preferably 10°C/second or less.

In addition, from the viewpoint of operation, the average heating rate is set to 0.5° C./sec or more. The average heating rate is preferably 1.0° C./sec or more, and more preferably 1.5° C./sec or more.

Here, the average heating rate (°C/second) is calculated from ((Ac1(°C)-500°C)/(heating time from 500°C to Ac1(°C) (seconds)).

Heat to the annealing temperature T below 840°C and satisfying 0.5≤(T-Ac1)/(Ac3-Ac1)<1.0

After the heating, the annealing temperature T is maintained in a furnace atmosphere with a dew point Td of -50°C or higher and -30°C or lower.

In the present invention, by heating to the annealing temperature T described later or further holding at the annealing temperature T, the needle-shaped γ described later can be formed. In this regard, if heated to the γ single-phase region above Ac3 (°C), the needle-shaped γ and the adjacent γ are combined, and the morphology of γ becomes equiaxed. Therefore, in the present invention, it is necessary to implement dual-phase region annealing.

In addition, regarding the annealing temperature T, if (T-Ac1)/(Ac3-Ac1) is less than 0.5, the reverse phase transformation to γ does not occur sufficiently, and needle-shaped γ is not formed, and only equiaxed γ is formed along the recrystallized ferrite grain boundary. Therefore, regarding the annealing temperature T, it is set to 0.5≤(T-Ac1)/(Ac3-Ac1)<1.0.

When the temperature T is higher than 840° C., good chemical conversion treatability cannot be obtained. Therefore, the temperature T is set to 840° C. or lower.

In addition, when the dew point Td is lower than -50°C, the formation of Si oxides that have an adverse effect on chemical conversion treatability is promoted, so good chemical conversion treatability cannot be obtained. In addition, when the dew point Td exceeds -30°C, an internal oxide layer of oxides is selectively formed at the grain boundaries generated in the structure, which has an adverse effect on corrosion resistance and the like. Therefore, the dew point Td is set to be higher than -50°C and lower than -30°C. The dew point Td is preferably higher than -48°C, and more preferably higher than -46°C. In addition, the dew point Td is preferably lower than -32°C, and more preferably lower than -34°C.

The soaking time at the annealing temperature T is not particularly limited, but is preferably set to 25 to 350 seconds, more preferably 50 to 300 seconds, from the viewpoint of element distribution in the two-phase region annealing.

In addition, Ac1 (° C.) can be calculated by the following formula based on the empirical rule.

Ac1(℃)＝723+22×[Si%]-18×[Mn%]+17×[Cr%]+4.5[Mo%]+16×[V%]

Ac3 (°C) can be calculated by the following formula based on the empirical rule.

Ac3(℃)=910-203×[C%] ^1/2 +44.7×[Si%]-30×[Mn%]+700×[P%]+400×[sol.Al%]-20×[ Cu%]+31.5×[Mo%]+104×[V%]+400×[Ti%]

In addition, [X%] in the above formula is the content (mass %) of the component element X in the steel plate, and is sometimes set to "0" when it is not contained.

The needle-shaped γ structure formed by the above-mentioned heat-maintaining treatment has a number density of 5 pieces/1000 μm ² or more.

In the present invention, in order to have the desired formability, needle-shaped γ is used. If a large amount of needle-shaped austenite (needle-shaped γ) is formed, a large amount of residual γ with a high aspect ratio is easily formed. In the present invention, in order to have the desired formability, it is necessary to make the needle-shaped γ structure formed by heating to the annealing temperature T and isothermal holding have a number density of 5 pieces/1000μm2 or ^more . Considering the properties of needle-shaped γ, there is no upper limit, and the number of needle-shaped γ particles is preferably large.

In the present invention, for a cold-rolled steel sheet having the above-mentioned component composition and structure, the average heating rate in the temperature range of 500°C to Ac1 is set to 0.5 to 15°C/second or less, and the steel sheet is heated to an annealing temperature T, and is heat-maintained at the annealing temperature T in a furnace atmosphere at a dew point Td, thereby adjusting the number density of the needle-shaped γ structure to a desired range.

Number density of needle-like γ structures

When evaluating a structure formed at a high temperature, the structure formed by freezing the structure by water cooling is usually evaluated. In the present invention, it is important that the needle-shaped γ formed by the treatment of isothermal holding until the annealing temperature T in the annealing process contributes to the formation of residual γ with a high aspect ratio and high processing stability in the subsequent cooling process, and the number density of the needle-shaped γ structure is measured. An observation sample is cut from the steel plate in a manner such that the cross section parallel to the rolling direction is used as the observation surface, and the plate thickness cross section is corroded and revealed with a 1 volume % nitric acid ethanol solution. The scanning electron microscope (SEM) is magnified to 2000 times, and a structure photograph is taken in an area of 3000μm2 or ^more at the plate thickness t/4. The SEM photograph shown in Figure 1 (b) is a photograph of the structure that is water-cooled after being kept at a temperature T within the scope of the present invention in the annealing process, forming needle-shaped γ and blocky γ, ferrite structure. Figure 2 shows a schematic diagram of the method for determining the aspect ratio of needle-shaped γ. Here, austenite with an aspect ratio of 3.0 or more surrounded by recrystallized ferrite having the same orientation is defined as needle-shaped γ. It should be noted that the tip of the needle-shaped austenite does not matter even if it contacts other austenite grains. In this case, it is sufficient to confirm that the adjacent ferrite grains are in the same orientation by electron beam backscatter diffraction (EBSD). According to this definition, in the annealing process, the number of needle-shaped γ in the steel sheet subjected to the soaking treatment to the annealing temperature T is measured in 5 fields of view, and the number density (number/1000 μm ² ) of needle-shaped γ is determined by dividing the number of needle-shaped γ by the total area of observation.

First cooling: Set the average cooling rate in the temperature range of 750-550°C to 6.0°C/s or more, and cool to the first cooling stop temperature Tc1 below 550°C and above 400°C First holding: After the first cooling, hold at the first cooling stop temperature Tc1 for more than 25 seconds

In the first cooling, ferrite transformation and pearlite transformation mainly occur in the temperature range of 750 to 550°C. If ferrite transformation or pearlite transformation occurs excessively, the needle-shaped γ undergoes ferrite transformation. Therefore, the average cooling rate in the temperature range of 750 to 550°C is set to 6.0°C/sec or more to suppress ferrite transformation. The average cooling rate is preferably 8.0°C/sec or more, and more preferably 10.0°C/sec or more.

Here, the average cooling rate (° C./second) is calculated from (750° C. (cooling start temperature)−550° C. (cooling stop temperature))/(cooling time (seconds) from the cooling start temperature to the cooling stop temperature).

The first cooling stop temperature Tc1 in the first cooling corresponds to the temperature on the high temperature side of the two-stage austenite tempering, and is set to the temperature range of upper bainite transformation with less carbide precipitation, that is, 550°C or less and 400°C or more. When the first cooling stop temperature Tc1 exceeds 550°C, ferrite transformation or pearlite transformation occurs, local carbon distribution does not occur, and residual γ with high processing stability is not formed. On the other hand, when the first cooling stop temperature Tc1 is lower than 400°C, lower bainite transformation accompanied by carbide precipitation occurs. In addition, depending on the component composition and annealing conditions, a part of the structure undergoes martensitic transformation and becomes tempered martensite through subsequent maintenance. As a result, the carbon distribution to the untransformed γ is delayed, and residual γ with high processing stability is not formed. Therefore, the first cooling stop temperature Tc1 is set to be 550°C or less and 400°C or more. The first cooling stop temperature Tc1 is preferably 420°C or more, and more preferably 450°C or more. In addition, the first cooling stop temperature Tc1 is preferably 530°C or lower, and more preferably 510°C or lower.

In addition, as the first holding after the first cooling, the temperature control at the first cooling stop temperature Tc1 is allowed as long as it is within the range of 550°C or less and 400°C or more, and is held for more than 25 seconds. In this way, the upper bainite can be fully transformed. The holding time in the first holding is preferably 30 seconds or more, and more preferably 35 seconds or more. In addition, the holding time in the first holding is preferably 60 seconds or less, and more preferably 55 seconds or less.

Second cooling: Cooling to a second cooling stop temperature Tc2 which is lower than the first cooling stop temperature Tc1 and lower than 450°C and higher than 300°C

Second hold: hold at the second cooling stop temperature Tc2 for 20 to 3000 seconds

In the second cooling, the second cooling stop temperature Tc2 is equivalent to the temperature on the low temperature side of the two-stage austenite tempering. By maintaining the steel plate at the second cooling stop temperature Tc2, the amount of solid solution C that undergoes martensitic transformation during the final cooling after the first cooling stop temperature Tc1 is maintained is small, so that the untransformed γ with low thermal stability undergoes bainite transformation. Thus, residual γ can be formed. When the second cooling stop temperature Tc2 is a temperature below the first cooling stop temperature Tc1 and exceeds 450°C, the bainite transformation stops, so the untransformed γ undergoes martensitic transformation excessively during the final cooling. On the other hand, when the second cooling stop temperature Tc2 is lower than 300°C, the untransformed γ undergoes martensitic transformation, and carbide formation is promoted by maintaining. Thus, the formation of residual γ with a high amount of solid solution C and high processing stability is hindered. Therefore, the second cooling stop temperature Tc2 is set to be below Tc1 and below 450°C and above 300°C. The second cooling stop temperature Tc2 is preferably above 320°C, and more preferably above 340°C. In addition, the second cooling stop temperature Tc2 is preferably 430°C or lower, and more preferably 410°C or lower.

In addition, as the second holding after the second cooling, the temperature control at the second stop temperature Tc2 is allowed as long as it is below the first cooling stop temperature Tc1 and below 450°C and above 300°C. When the holding time is 20 seconds or more, the formation of residual γ with high processing stability is promoted. On the other hand, the holding time in the second holding is set to 3000 seconds or less from the operational point of view. The holding time in the second holding is preferably 100 seconds or more, and more preferably 200 seconds or more. In addition, the holding time in the second holding is preferably 2500 seconds or less, and more preferably 2000 seconds or less.

Third cooling: After the second hold, cooling

After the second holding, the steel sheet is cooled to room temperature (10 to 30° C.) to obtain the steel sheet of the present invention.

The process after the annealing step is not particularly limited, and for example, temper rolling with an elongation of 0.05 to 0.5% may be performed.

The steel sheet of the present invention obtained by the method for producing a steel sheet of the present invention preferably has a thickness of 0.5 mm or more and preferably has a thickness of 2.0 mm or less.

Next, the member of the present invention and a method for producing the same will be described.

The component of the present invention is formed by subjecting the steel plate of the present invention to at least one of forming and joining. In addition, the method for manufacturing the component of the present invention includes the step of subjecting the steel plate of the present invention to at least one of forming and joining to form the component.

The steel sheet of the present invention has a tensile strength of 590 MPa or more, and has high ductility, excellent stretch flange formability, and good chemical conversion treatment properties. Therefore, the components obtained using the steel sheet of the present invention are also high-strength, and have excellent high ductility, excellent stretch flange formability, and good chemical conversion treatment properties compared to conventional high-strength components. In addition, if the components of the present invention are used, lightweighting can be achieved. Therefore, the components of the present invention can be appropriately used, for example, for vehicle body frame parts. The components of the present invention also include welded joints.

The forming process can use general processing methods such as press working without limitation. In addition, the joining process can use general welding such as spot welding and arc welding, rivet joining, caulking joining, etc. without limitation.

Example

Hereinafter, embodiments of the present invention will be described.

A steel slab with a thickness of 250 mm and a component composition shown in Table 1 was subjected to hot rolling (slab heating temperature: 1250° C., soaking time: 30 minutes, finishing rolling temperature: Ar3+50° C., coiling temperature: 550° C.) and pickling, and the obtained hot-rolled steel sheet was cold-rolled under the conditions shown in Table 2 to produce a cold-rolled steel sheet.

Next, the cold-rolled steel sheets were annealed in a continuous annealing line under the conditions shown in Table 2, and then temper rolled at an elongation of 0.2 to 0.4% to produce steel sheets for evaluation.

In addition, after the second holding at the second cooling stop temperature Tc2, the temperature is cooled to room temperature (20°C) as the third cooling.

The obtained steel sheets were evaluated by the following methods.

(1) Determination of the area ratio of steel structure and the number density of needle-shaped γ

The observation sample was cut from the steel plate in a manner that the cross section perpendicular to the steel plate surface and parallel to the rolling direction was used as the observation surface, and the plate thickness cross section was corroded with a 1 volume % nitric acid ethanol solution, and the scanning electron microscope (SEM) was magnified to 2000 times, and the organization photograph was taken in the area of 3000 μm ² or more at the plate thickness t/4. The following items (i) to (iv) were measured respectively. The results are shown in Table 3. It should be noted that t represents the plate thickness and w represents the plate width.

(i) Polygonal ferrite and upper bainite

Polygonal ferrite (recrystallized F) and upper bainite (UB) both appear gray in the SEM photograph, but can be identified by their shapes. An example of an SEM photograph is shown in FIG1 together with an SEM photograph of the structure after water cooling after being held at temperature T. The area represented by the dotted line in FIG1(a) is a needle-shaped γ structure formed by a treatment of isothermal holding at an annealing temperature T within the scope of the present invention in the annealing process, in which upper bainite (UB) is generated, and residual γ or fresh martensite (M) with a high aspect ratio is formed around it. The same structure is also observed in the blocky γ structure formed by isothermal holding at an annealing temperature T. The area ratios of polygonal ferrite and upper bainite are measured by the point algorithm according to ASTM E562-11 (2014). The area ratio of polygonal ferrite and the area ratio of upper bainite are the average values of the measured values at 5 locations, respectively.

(ii) Fresh martensite and residual γ

Both fresh martensite and residual γ appear white in the SEM photograph and cannot be distinguished. Therefore, residual γ is measured separately by the method described later. In addition, the total area ratio of fresh martensite and residual γ is measured by the point algorithm according to ASTME562-11 (2014) from the SEM photograph, and the area ratio of residual γ measured by the method described later is subtracted from the total area ratio to determine the area ratio of fresh martensite. The total area ratio of fresh martensite and residual γ is measured by the point algorithm, and the value obtained by subtracting the volume ratio of residual γ measured by the method described later from the average value of the measured values at 5 places is set as the area ratio of fresh martensite.

(iii) Tempered martensite and/or lower bainite

Tempered martensite and lower bainite are structures containing carbides observed as white fine structures in SEM photos. The two can be distinguished in a more microscopic observation, but it is difficult to distinguish in SEM photos. Therefore, in the present invention, tempered martensite and lower bainite are defined as the same structure, and the total area ratio of tempered martensite and lower bainite is measured by point calculation according to ASTM E562-11 (2014). The average value of the measured values at 5 locations is used as the total area ratio of tempered martensite and lower bainite.

(iv) Surplus organization

The area ratio of the remaining structure is defined as the area ratio of the remaining structure by subtracting the area ratios of polygonal ferrite, upper bainite, fresh martensite, retained γ, tempered martensite and lower bainite measured by the above method from 100%.

(2) Determination of volume fraction of retained γ

After the steel plate is ground to a position of 1/4 of the plate thickness, it is further ground by 0.1 mm by chemical grinding. For the surface thus obtained, the integrated reflection intensity of the (200) plane, (220) plane, (311) plane of FCC iron (γ) and the (200) plane, (211) plane, (220) plane of BCC iron (ferrite) is measured by using Mo Kα rays using an X-ray diffraction device, and the volume ratio of residual γ is determined from the intensity ratio of the integrated reflection intensity from each surface of FCC iron (γ) to the integrated reflection intensity from each surface of BCC iron (ferrite). In the present invention, the volume ratio of the residual γ can be used as the area ratio of the residual γ.

(3) Number density of needle-shaped γ structures

When evaluating a structure formed at a high temperature, the structure formed by freezing the structure by water cooling is usually evaluated. In the present invention, it is important that the needle-shaped γ formed by the treatment of isothermal holding until the annealing temperature T in the annealing process contributes to the formation of residual γ with a high aspect ratio and high processing stability in the subsequent cooling process, and the number density of the needle-shaped γ structure is measured. An observation sample is cut from the steel plate in a manner that the cross section parallel to the rolling direction is used as the observation surface, and the plate thickness cross section is corroded and revealed with a 1 volume % nitric acid ethanol solution. The scanning electron microscope (SEM) is magnified to 2000 times, and a structure photograph is taken in an area of 3000μm2 or ^more at the plate thickness t/4. The SEM photograph shown in Figure 1 (b) is a photograph of the structure that is water-cooled after being kept at a temperature T within the scope of the present invention in the annealing process, forming needle-shaped γ and blocky γ, ferrite structure. Figure 2 shows a schematic diagram of the method for determining the aspect ratio of needle-shaped γ. Here, austenite with an aspect ratio of 3.0 or more surrounded by recrystallized ferrite having the same orientation is defined as needle-shaped γ. It should be noted that the tip of the needle-shaped austenite does not matter even if it contacts other austenite grains. In this case, it is sufficient to confirm that the adjacent ferrite grains are in the same orientation by electron beam backscatter diffraction (EBSD). According to this definition, in the annealing process, the number of needle-shaped γ in the steel plate subjected to the soaking treatment to the annealing temperature T is measured in 5 fields of view, and the number density (number/ ^1000μm2 ) of needle-shaped γ is determined by dividing the number of needle-shaped γ by the total area observed. The results are shown in Table 3.

(4) Equivalent circular diameter and aspect ratio of fresh martensitic particles and/or residual γ particles

An observation sample was cut from the steel plate with the cross section parallel to the rolling direction as the observation surface, and the structural corrosion of the plate thickness cross section was visualized using Lepera etching solution. A laser microscope (LM) was used to magnify the area to 1000 times at a thickness of t/4 and above 10000 ^μm2 to take a structural photograph.

Lepera etching is a color etching method that represents fresh martensite and/or residual γ with a white contrast, extracts fresh martensite particles and/or residual γ particles, and performs image analysis to determine the equivalent circular diameter and aspect ratio of the fresh martensite particles and/or residual γ particles.

Among all the obtained particles, the number of particles having an equivalent circle diameter of less than 1.2 μm was measured, and the ratio of the number of particles to the number of all the particles was calculated.

In addition, among all the obtained particles, the number of particles with an aspect ratio of 2.5 or more was measured, with particles with an equivalent circle diameter of 1.2 μm or more as the object, and the ratio of particles with an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more to all particles with an equivalent circle diameter of 1.2 μm or more was calculated. The results are shown in Table 3.

(5) Texture of cold rolled structure

A measurement sample having a cross section parallel to the rolling direction as a measurement surface is cut from the cold-rolled steel sheet after the cold rolling process, and the measurement surface is mechanically polished or electrolytically polished, and then the area of 80000 μm ² or more is measured by the SEM-EBSD method (measurement conditions: WD: 20 mm, acceleration voltage: 20 kV). The area ratio of the bcc phase with the {ND plane} <RD direction> orientation of the rolling {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100} <011> orientation is quantified, and the ratio to the area ratio of the bcc phase of all orientations is calculated, thereby evaluating the texture of the cold-rolled steel sheet.

(6) Tensile test

From the obtained steel plate, a JIS No. 5 tensile test piece having a tensile direction in a direction perpendicular to the rolling direction was prepared. Each test piece was subjected to a tensile test in accordance with the provisions of JIS Z 2241 (2011). The crosshead speed of the tensile test was set to 10 mm/min. It should be noted that the measurement was performed twice, and the measured value was averaged to obtain the tensile strength (TS) of each steel plate.

(7) Hole expansion test

A test piece of 100 mm×100 mm was cut out, and a hole expansion test based on JFST 1001 (Japan Iron and Steel Federation Standard) was performed three times at each cutting position, and the average value of the three times (total value (%) of the three times/3) was taken as the hole expansion ratio λ (%).

(8) Evaluation

In the present invention, a steel sheet having a tensile strength (TS) of 590 MPa or more is evaluated as having high strength.

When the tensile strength (TS) × total elongation (T.El) ≥ 22000 MPa·% or above, the ductility El is evaluated as excellent, and when the hole expansion ratio λ (%) satisfies the following (A1) or (A2), the stretched flange formability λ is evaluated as excellent.

(A1) When the tensile strength is 590 MPa or more and less than 780 MPa, the hole expansion ratio λ is 60% or more

(A2) When the tensile strength is 780 MPa or more, the hole expansion ratio λ is 35% or more

(9) Chemical conversion treatability

The annealed steel sheet was electrolytically pickled with sulfuric acid at a current density of 20 to 35 A/ ^dm2 for 2 seconds, degreased, and surface adjusted, and then chemically converted using a zinc phosphate chemical conversion treatment solution. Degreasing process: treatment temperature 40°C, treatment time 120 seconds, spray degreasing; surface adjustment process: pH 9.5, treatment temperature room temperature, treatment time 20 seconds; chemical conversion treatment process: chemical conversion treatment was performed under the conditions of a temperature of the chemical conversion treatment solution of 35°C and a treatment time of 120 seconds. It should be noted that as treatment agents in each process of the degreasing process, surface adjustment process, and chemical conversion treatment process, the degreasing agent FC-E2011, the surface conditioning agent PL-X, and the chemical conversion treatment solution パルボンドPB-L3065 manufactured by Japan Parkerizing Co., Ltd. were used in sequence. The surface chemical conversion structure was observed by SEM observation at a magnification of 2000 times in an area of 10000 μm ² or more, and the sample in which the chemical conversion film structure was observed on the entire surface was evaluated as 0, and the sample in which the chemical conversion film structure was not formed partially by visual observation was evaluated as ×. The results are shown in Table 3.

As shown in Table 3, it can be seen that the steel sheet of the present invention has a tensile strength of 590 MPa or more, has high ductility and excellent stretch flange formability, and is also excellent in chemical conversion treatability.

It can also be seen that since the steel plate of the example of the present invention is high strength, has high ductility, excellent stretch flange formability and good chemical conversion treatment property, the components obtained by forming the steel plate of the example of the present invention, the components obtained by joining, and the components obtained by forming and joining are high strength, high ductility, excellent stretch flange formability and good chemical conversion treatment property like the steel plate of the example of the present invention.

Claims (6)

1. A steel sheet having a composition comprising, by mass%, 0.06 to 0.24% C, 0.4% or more and less than 1.60% Si, 1.5 to 3.2% Mn, 0.02% or less P, 0.01% or less S, less than 1.0% sol.Al, less than 0.015% N, and satisfying the following formula (1), with the balance being Fe and unavoidable impurities, Si/Mn<0.50…Formula (1) Here, in formula (1), Si and Mn represent Si content (mass %) and Mn content (mass %), respectively. The steel plate has a structure in which the area ratio of polygonal ferrite is 20% to 85%, the area ratio of upper bainite is 9% to 45%, the volume ratio of retained austenite is 3% to 15%, the area ratio of fresh martensite is 3% to 15%, the total area ratio of tempered martensite and lower bainite is 50% or less (including 0%), and the area ratio of residual structure is 5% or less, The ratio of the total number of fresh martensite particles and retained austenite particles having an equivalent circle diameter of less than 1.2 μm to the total number of fresh martensite particles and retained austenite particles is 50% or more, Furthermore, the ratio of the total number of fresh martensite particles and retained austenite particles having an aspect ratio of 2.5 or more and an equivalent circle diameter of 1.2 μm or more to the number of fresh martensite particles and retained austenite particles having an equivalent circle diameter of 1.2 μm or more is 40% or more. 2. The steel sheet according to claim 1, wherein the component composition further contains, in mass %, one or more selected from the group consisting of Nb: 0.2% or less, Ti: 0.2% or less, V: 0.2% or less, B: 0.01% or less, Cu: 0.2% or less, Ni: 0.2% or less, Cr: 0.4% or less, and Mo: 0.15% or less. 3. The steel sheet according to claim 1 or 2, further comprising, as the component composition, one or more selected from the group consisting of Mg: 0.0050% or less, Ca: 0.0050% or less, Sn: 0.10% or less, Sb: 0.10% or less, and REM: 0.0050% or less in mass %. 4 . A member formed using the steel plate according to claim 1 . 5. A method for manufacturing a steel plate, comprising: a cold rolling step of hot-rolling and pickling the steel slab having the composition of any one of claims 1 to 3, and then cold-rolling the obtained hot-rolled steel sheet to obtain a cold-rolled steel sheet; and an annealing step of subjecting the cold-rolled steel sheet to an annealing treatment to obtain a steel sheet; The cold rolling process includes the following steps: setting the cumulative cold rolling ratio to 30 to 85%, setting the first pass reduction ratio to 5% or more and less than 25%, thereby obtaining the cold rolled steel sheet having a total of structures having a {111}<0-11> orientation, a {111}<11-2> orientation, a {211}<0-11> orientation and a {100}<011> orientation of 35% or more and 75% or less in terms of area ratio relative to the total structure of the bcc phase, The annealing process includes the following annealing treatment: The cold-rolled steel sheet is heated to an annealing temperature T of 840° C. or less and satisfying 0.5≤(T-Ac1)/(Ac3-Ac1)<1.0, with an average heating rate of 0.5 to 15° C./second in a temperature range of 500° C. or more and Ac1 or less. After the heating, the steel sheet is held at the annealing temperature T in a furnace atmosphere having a dew point Td of -50°C or higher and -30°C or lower, thereby obtaining a steel sheet having a needle-like austenite structure with a number density of 5 or more per 1000 μm ² . Next, the first cooling is performed by setting the average cooling rate in the temperature range of 750 to 550° C. to 6.0° C./sec or more, and cooling to a first cooling stop temperature Tc1 of 550° C. or less and 400° C. or more. After the first cooling, a first holding is performed at the first cooling stop temperature Tc1 for more than 25 seconds. After the first holding, a second cooling is performed to cool to a second cooling stop temperature Tc2 which is lower than the first cooling stop temperature Tc1 and lower than 450° C. and higher than 300° C. A second holding period of 20 to 3000 seconds is performed at the second cooling stop temperature Tc2. After the second holding, a third cooling step of cooling is performed. 6 . A method for manufacturing a member, comprising the step of subjecting the steel sheet according to claim 1 to at least one of a forming process and a joining process to produce a member.