JP7674676B2 - Steel plate and its manufacturing method - Google Patents

Description

This application discloses a steel plate and a method for manufacturing the same.

In recent years, efforts have been made to reduce the weight of automobile bodies by using high-strength steel plates in order to improve fuel efficiency. Also, to ensure the safety of passengers, high-strength steel plates are increasingly being used in place of mild steel plates in automobile bodies. In order to further reduce the weight of automobile bodies in the future, the strength level of high-strength steel plates must be increased even more than before.

In addition, automotive parts are required to exhibit the function of suppressing deformation during an automobile collision. In order to increase the resistance of automotive parts to deformation during an automobile collision, it is desirable to increase the bending strength of the automotive parts. In addition, in order to increase the bending strength from a structural standpoint by optimizing the part shape, high formability is required for the steel plate. For this reason, the steel plate used in automotive parts is required to have high strength, excellent bending strength, and also exhibit high elongation. However, in the prior art, although the workability of high-strength steel plate has been studied (for example, Patent Documents 1 to 3 below), sufficient study has not been made on ensuring bending strength against bending deformation from both the front and back sides.

Patent Document 1 discloses a high-strength steel sheet with excellent workability, which has ferrite as the main phase, contains an average of 5 volume% or more of retained austenite, and has a difference ΔVγ between the maximum and minimum retained austenite content at each position in the sheet thickness direction between 0.1 mm from the surface of the steel sheet and 0.1 mm from the back surface of the steel sheet of 3.0 volume% or less.

Patent Document 2 discloses a steel plate for ship hulls with excellent impact absorbing ability that can minimize destruction of the hull in the event of a collision with a tanker. The steel plate contains C, Si, Mn, and Al, and further contains strengthening elements as necessary, with the balance being Fe and unavoidable impurities, and has a thickness of 8 mm or more, and is characterized in that the front and back layers of the steel plate, at least 1/8 of the thickness of the steel plate, contain residual γ at an area ratio of 1.0 to 20%.

Patent Document 3 discloses a thick structural steel plate that can dramatically improve both brittle crack propagation arrest properties and Charpy properties without relying on the addition of expensive alloy elements such as Ni. The steel plate is composed of, by weight, C: 0.04-0.30%, Si: ≦0.5%, Mn: ≦2.0%, Al: ≦0.1%, Ti: 0.001-0.10%, N: 0.001-0.01%, with the balance being Fe and unavoidable impurities, and in which the average crystal grain size d of the structure in specified regions of the front and back layers of the plate thickness is 3 μm or less, and the Vickers hardness of the structure satisfies specified requirements.

Patent No. 3546266 Patent No. 3499126 Patent No. 3845113

In view of the above-mentioned circumstances, the present application discloses a steel plate having excellent mechanical properties such as strength and elongation as well as excellent bending strength, and a manufacturing method thereof.

The inventors have conducted extensive research into methods for solving the above problems and have demonstrated that by optimizing the proportion of steel plate structures, including retained austenite, and reducing the difference in number density of precipitates on the front and back surfaces, it is possible to obtain a steel plate that has excellent mechanical properties such as strength and elongation as well as high bending strength. Additionally, it has been confirmed that in steel plates with a difference in number density of precipitates on the front and back surfaces of more than 10%, the bending strength changes depending on the bending direction, and the deformation resistance of parts during a collision is accidentally reduced.

The inventors have also discovered that an integrated manufacturing method, which involves subjecting a cold-rolled sheet to two annealing processes, and then coiling and recoiling the sheet between the two annealing processes to a specified aging treatment, can produce a steel sheet with an optimized structure and a small difference in the number density of precipitates on the front and back surfaces.

Furthermore, through accumulating various research studies, the inventors have also discovered that steel plate in which bending strength is increased by reducing the difference in number density of precipitates on the front and back surfaces as described above is difficult to manufacture by simply devising a single set of hot rolling conditions or annealing conditions, and can only be manufactured by achieving optimization in a so-called integrated process, including the hot rolling and annealing processes.

The gist of the present invention is as follows.
(1)
In mass percent,
C: 0.10-0.30%,
Si: 0.60-1.20%,
Mn: 1.00-3.50%,
P: 0.0200% or less,
S: 0.0200% or less,
Al: 0.001-1.000%,
N: 0.0200% or less,
Ti: 0 to 0.500%,
Co: 0 to 0.500%,
Ni: 0 to 0.500%,
Mo: 0-0.500%,
Cr: 0-2.000%,
O: 0 to 0.0100%,
B: 0 to 0.0100%,
Nb: 0 to 0.500%,
V: 0 to 0.500%,
Cu: 0-0.500%,
W: 0-0.1000%,
Ta: 0-0.1000%,
Sn: 0-0.0500%,
Sb: 0 to 0.0500%,
As: 0 to 0.0500%,
Mg: 0 to 0.0500%,
Ca: 0-0.0500%,
Y: 0 to 0.0500%,
Zr: 0 to 0.0500%,
La: 0 to 0.0500%, and
Ce: 0 to 0.0500%,
and the balance being Fe and impurities,
In terms of area ratio,
The sum of ferrite, pearlite and bainite: 0% or more and 30.0% or less, and
Retained austenite: 10.0% or more and 30.0% or less,
The balance has a steel structure consisting of fresh martensite and tempered martensite,
a difference between the number density of precipitates in the tempered martensite in the first surface on the front side of the steel sheet and the number density of precipitates in the tempered martensite in the second surface on the back side of the steel sheet is within 10.0%;
The yield strength is 600 MPa or more.
Steel plate.
(2)
In mass percent,
Ti: 0.001 to 0.500%,
Co: 0.001 to 0.500%,
Ni: 0.001 to 0.500%,
Mo: 0.001-0.500%,
Cr:0.001~2.000%
O: 0.0001-0.0100%
B: 0.0001 to 0.0100%,
Nb: 0.001-0.500%,
V: 0.001-0.500%,
Cu: 0.001 to 0.500%,
W: 0.0001-0.1000%,
Ta: 0.0001 to 0.1000%,
Sn: 0.0001 to 0.0500%,
Sb: 0.0001 to 0.0500%,
As: 0.0001 to 0.0500%,
Mg: 0.0001-0.0500%,
Ca: 0.0001-0.0500%,
Y: 0.0001-0.0500%,
Zr: 0.0001 to 0.0500%,
La: 0.0001 to 0.0500%, and Ce: 0.0001 to 0.0500%,
The chemical composition contains one or more of the following:
The steel plate according to (1) above.
(3)
The steel structure contains the acicular retained austenite.
The steel sheet according to (1) or (2) above.
(4)
A method for manufacturing a steel sheet, comprising the steps of:
Hot rolling a steel slab having the chemical composition described in (1) or (2) above to obtain a hot-rolled sheet;
coiling the hot-rolled sheet;
pickling the hot-rolled sheet;
cold rolling the hot rolled sheet to obtain a cold rolled sheet;
Q annealing (Q: Quenching) is performed on the cold-rolled sheet;
The cold-rolled sheet subjected to the Q annealing is subjected to IA annealing (IA: Intercritical Annealing); and
performing an aging treatment between the Q annealing and the IA annealing;
Including,
The Q annealing is a process of heating the cold-rolled sheet to a temperature in the austenite single phase region and 1000° C. or less, and then cooling the sheet to obtain a martensite structure having an area ratio of 90.0% or more.
The IA annealing is a process of holding the cold-rolled sheet in a two-phase region of ferrite and austenite to obtain retained austenite,
The aging treatment includes aging treatment 1 in which one of the front and back sides of the cold-rolled sheet is subjected to tensile deformation with a bending R of 2.0 m or less and held at 0 to 40 ° C. for 20 hr or more, and aging treatment 2 in which the other of the front and back sides of the cold-rolled sheet is subjected to tensile deformation with a bending R of 2.0 m or less and held at 0 to 40 ° C. for 20 hr or more.
Manufacturing method of steel plate.
(5)
In the IA annealing, the cold-rolled sheet is held in a two-phase region of ferrite and austenite, and then cooled to room temperature. In the process, a coating layer made of zinc, aluminum, magnesium, or an alloy thereof is formed on the front and back surfaces of the cold-rolled sheet.
The manufacturing method described in (4) above.
(6)
In the IA annealing, needle-shaped retained austenite is obtained.
The manufacturing method according to (4) or (5).

The steel plate disclosed herein has excellent mechanical properties such as strength and elongation, as well as excellent bending strength.

Hereinafter, embodiments of the present invention will be described. Note that these descriptions are intended to be merely examples of embodiments of the present invention, and the present invention is not limited to the following embodiments.

<Steel Plate>
The steel plate according to this embodiment has, in mass%,
C: 0.10-0.30%,
Si: 0.60-1.20%,
Mn: 1.00-3.50%,
P: 0.0200% or less,
S: 0.0200% or less,
Al: 0.001-1.000%,
N: 0.0200% or less,
Ti: 0 to 0.500%,
Co: 0 to 0.500%,
Ni: 0 to 0.500%,
Mo: 0-0.500%,
Cr: 0-2.000%,
O: 0 to 0.0100%,
B: 0 to 0.0100%,
Nb: 0 to 0.500%,
V: 0 to 0.500%,
Cu: 0-0.500%,
W: 0-0.1000%,
Ta: 0-0.1000%,
Sn: 0-0.0500%,
Sb: 0 to 0.0500%,
As: 0 to 0.0500%,
Mg: 0 to 0.0500%,
Ca: 0-0.0500%,
Y: 0 to 0.0500%,
Zr: 0 to 0.0500%,
La: 0 to 0.0500%, and
Ce: 0 to 0.0500%,
and the balance being Fe and impurities,
In terms of area ratio,
The sum of ferrite, pearlite and bainite: 0% or more and 30.0% or less, and
Retained austenite: 10.0% or more and 30.0% or less,
The balance has a steel structure consisting of fresh martensite and tempered martensite,
a difference between the number density of precipitates in the tempered martensite in the first surface on the front side of the steel sheet and the number density of precipitates in the tempered martensite in the second surface on the back side of the steel sheet is within 10.0%;
It is characterized by a yield strength of 600 MPa or more.

First, we will explain the reasons for limiting the chemical composition of the steel plate according to the embodiment of the present invention. Here, "%" for the components means mass %. Furthermore, in this specification, unless otherwise specified, "to" indicating a numerical range is used to mean that the numerical values written before and after it are included as the lower and upper limits.

(C: 0.10-0.30%)
C is an element that inexpensively increases tensile strength and is an extremely important element for controlling the strength of steel. When the C content is 0.10% or more, such an effect is easily obtained. The C content may be 0.12% or more. On the other hand, excessive C content may lead to a decrease in elongation and brittle fracture of the steel, which may promote a decrease in bending strength during deformation of the part. When the C content is 0.30% or less, such problems are easily avoided. The C content may be 0.28% or less.

(Si: 0.60-1.20%)
Si acts as a deoxidizer, and is an element that increases the stability of the retained austenite structure against processing and suppresses the precipitation of carbides in the martensite structure during aging. When the Si content is 0.60% or more, such effects are easily obtained. The Si content may be 0.70% or more. On the other hand, when Si is contained excessively, the generation of ε carbides is suppressed during aging treatment, and bending strength may be reduced. When the Si content is 1.20% or less, such problems are easily avoided. The Si content may be 1.00% or less.

(Mn: 1.00-3.50%)
Mn is a factor that affects the ferrite transformation of steel, and is an element that suppresses the ferrite transformation in the cooling process of Q annealing described later, increases the structure ratio of martensite after Q annealing, and is effective in increasing strength. When the Mn content is 1.00% or more, such effects are easily obtained. The Mn content may be 1.30% or more. On the other hand, when Mn is excessively contained, a Mn-enriched layer due to microsegregation and center segregation appears prominently in the steel sheet, and due to the difference in solidification rate on the front and back surfaces of the slab, there is a difference in the distribution state of the Mn-enriched layer on the front and back surfaces of the steel sheet, so that the difference in the formation of the Mn segregation band may lead to a difference in the bending strength on the front and back surfaces. When the Mn content is 3.50% or less, such problems are easily avoided. The Mn content may be 3.00% or less.

(P: 0.0200% or less)
P is an element that strongly segregates at ferrite grain boundaries and promotes embrittlement of the grain boundaries, and the less the better. In addition, excessive P content may cause brittle fracture of the steel and promote a decrease in bending strength during deformation of the part. In this respect, the P content is 0.0200% or less. The P content may be 0.0180% or less. On the other hand, the lower limit of the P content is not particularly limited. The P content is 0% or more, may be 0.0001% or more, or may be 0.0010% or more.

(S: 0.0200% or less)
S is an element that generates nonmetallic inclusions such as MnS in steel and reduces the ductility of steel parts, and the less the better. In addition, excessive S content may cause voids to be generated starting from the nonmetallic inclusions during deformation of the part, and may also reduce bending strength. In this regard, the S content is 0.0200% or less. The S content may be 0.0180% or less. On the other hand, the lower limit of the S content is not particularly limited. The S content is 0% or more, and may be 0.0001% or more, or may be 0.0005% or more.

(Al: 0.001-1.000%)
Al is an element that acts as a deoxidizer for steel and stabilizes ferrite, and is added as necessary. When the Al content is 0.001% or more, such an effect is easily obtained. The Al content may be 0.010% or more. On the other hand, if an excessive amount of Al is contained, the ferrite transformation and bainite transformation during the cooling process in annealing may be excessively promoted, resulting in a decrease in the strength of the steel sheet. When the Al content is 1.000% or less, such a problem is easily avoided. The Al content may be 0.800% or less.

(N: 0.0200% or less)
N is an element that forms coarse nitrides in the steel sheet and reduces the workability of the steel sheet. In addition, N is an element that causes blowholes during welding. In addition, excessive N is combined with Al or Ti to generate a large amount of AlN or TiN, and these nitrides may become the starting point of void generation during part deformation, resulting in a decrease in bending strength. In this regard, the N content is 0.0200% or less. The N content may be 0.0160% or less. On the other hand, the lower limit of the N content is not particularly limited. The N content is 0% or more, may be 0.0001% or more, or may be 0.0010% or more.

The basic chemical composition of the steel sheet in this embodiment is as described above. Furthermore, the steel sheet in this embodiment may contain at least one of the following optional elements as necessary. These elements do not have to be contained, so the lower limit is 0%.

(Ti: 0-0.500%)
Ti is a strengthening element. It contributes to increasing the strength of the steel plate through precipitation strengthening, fine grain strengthening by suppressing the growth of crystal grains, and dislocation strengthening through suppression of recrystallization. On the other hand, excessive Ti content increases the precipitation of coarse carbides, which may become the origin of void generation during deformation of parts, resulting in a decrease in bending strength. The Ti content is 0% or more, may be 0.001% or more, may be 0.005% or more, and may be 0.500% or less, or may be 0.400% or less.

(Co: 0-0.500%)
Co is an element effective for controlling the morphology of carbides and increasing strength, and is added as necessary to control strength. On the other hand, excessive Co content causes a large number of fine Co carbides to precipitate, and these carbides may become the starting point for void generation during deformation of the part, resulting in a decrease in bending strength. The Co content is 0% or more, and may be 0.001% or more, and may be 0.500% or less, or 0.400% or less.

(Ni: 0-0.500%)
Ni is a strengthening element and is effective in improving hardenability. In addition, Ni may be added because it improves the wettability between the steel sheet and the plating and promotes the alloying reaction. On the other hand, excessive Ni content may affect the peelability of oxide scale during hot rolling and promote the generation of scratches on the steel sheet surface, thereby reducing the yield strength during bending deformation. The Ni content is 0% or more, and may be 0.001% or more, and may be 0.500% or less, and may be 0.400% or less.

(Mo: 0-0.500%)
Mo is an element effective in improving the strength of steel sheets. Mo is also an element that has the effect of suppressing ferrite transformation that occurs during heat treatment in continuous annealing equipment or continuous hot-dip galvanizing equipment. On the other hand, if Mo is contained excessively, a large number of fine Mo carbides are precipitated, and these carbides may become the starting point of void generation during part deformation, resulting in a decrease in bending strength. The Mo content is 0% or more, and may be 0.001% or more, and may be 0.500% or less, or 0.400% or less.

(Cr: 0-2.000%)
Cr, like Mn, is an element that suppresses pearlite transformation and is effective in increasing the strength of steel, and is added as necessary. On the other hand, excessive Cr content promotes the formation of retained austenite, and the presence of excess retained austenite may lead to a decrease in bending strength. The Cr content is 0% or more, and may be 0.001% or more, and may be 2.000% or less, and may be 1.500% or less.

(O: 0-0.0100%)
Since O forms oxides and deteriorates workability, it is necessary to suppress the amount of O added. In particular, oxides often exist as inclusions, and when they exist on the punched end surface or the cut surface, they form notch-shaped scratches or coarse dimples on the end surface, which may lead to stress concentration during part deformation and may become the starting point for crack formation, resulting in a decrease in bending strength. The O content is 0.0100% or less, and may be 0.0080% or less. Note that the O content is 0% or more, but controlling the O content to less than 0.0001% may increase the refining time and increase the manufacturing cost. In order to prevent an increase in manufacturing costs, the O content may be 0.0001% or more, or may be 0.0010% or more.

(B: 0-0.0100%)
B is an element that suppresses the formation of ferrite and pearlite during the cooling process from austenite and promotes the formation of low-temperature transformation structures such as bainite or martensite. B is also an element that is beneficial for increasing the strength of steel and is added as necessary. On the other hand, excessive B content leads to the formation of coarse B inclusions in the steel, and these inclusions become the starting points for void generation, which may lead to a decrease in bending strength during part deformation. The B content is 0% or more, and may be 0.0001% or more, 0.0010% or more, and may be 0.0100% or less, or 0.0080% or less.

(Nb: 0-0.500%)
Nb is an element effective for controlling the morphology of carbides, and is also effective for improving toughness because its addition refines the structure. On the other hand, excessive Nb content causes a large number of fine and hard Nb carbides to precipitate, and these carbides become the starting points for void generation, which may lead to a decrease in bending strength during part deformation. The Nb content is 0% or more, and may be 0.001% or more, and may be 0.500% or less, or 0.400% or less.

(V: 0-0.500%)
V is a strengthening element. It contributes to increasing the strength of the steel sheet through precipitation strengthening, fine grain strengthening by suppressing the growth of ferrite crystal grains, and dislocation strengthening through suppression of recrystallization. On the other hand, excessive V content increases the precipitation of carbonitrides, and these carbonitrides become the starting point of void generation, which may lead to a decrease in bending strength during deformation of parts. The V content is 0% or more, and may be 0.001% or more, and may be 0.500% or less, or 0.400% or less.

(Cu: 0-0.500%)
Cu is an element effective in improving the strength of steel sheet. On the other hand, if Cu is contained excessively, the steel material becomes embrittled during hot rolling, making hot rolling difficult. Furthermore, since the ductility decreases with increasing strength of the steel, it may lead to a decrease in bending strength during deformation of the part. The Cu content is 0% or more, and may be 0.001% or more, and may be 0.500% or less, or 0.400% or less.

(W: 0-0.1000%)
W is effective in increasing the strength of steel sheets, and precipitates and crystallized materials containing W act as hydrogen trapping sites. On the other hand, excessive W content facilitates the progression of void generation starting from coarse carbides, which may lead to a decrease in bending strength during deformation of parts. The W content is 0% or more, and may be 0.0001% or more, 0.0010% or more, or 0.1000% or less, or 0.0800% or less.

(Ta: 0-0.1000%)
Ta, like Nb, V, and W, is an element effective in controlling the morphology of carbides and increasing strength, and is added as necessary. On the other hand, if Ta is contained excessively, a large number of fine Ta carbides are precipitated, and voids are easily generated starting from these carbides, which may lead to a decrease in bending strength during deformation of the part. The Ta content is 0% or more, and may be 0.0001% or more, 0.0010% or more, and may be 0.1000% or less, or 0.0800% or less.

(Sn: 0-0.0500%)
Sn is an element contained in steel when scrap is used as a raw material, and the less the better. Excessive Sn content may lead to a decrease in bending strength during part deformation due to embrittlement of the steel plate. The Sn content is 0.0500% or less, and may be 0.0400% or less. The Sn content may be 0%, but controlling the Sn content to less than 0.0001% may increase the refining time and increase the manufacturing cost. In order to prevent an increase in manufacturing costs, the Sn content may be 0.0001% or more, or may be 0.0010% or more.

(Sb: 0-0.0500%)
Sb, like Sn, is an element contained when scrap is used as a steel raw material. Sb strongly segregates at grain boundaries, which leads to embrittlement of the grain boundaries and a decrease in ductility, so the less Sb, the better. In addition, excessive Sb content may lead to a decrease in bending strength during part deformation due to embrittlement of the steel plate. The Sb content is 0.0500% or less, and may be 0.0400% or less. The Sb content may be 0%, but controlling the Sb content to less than 0.0001% may increase the refining time and increase the manufacturing cost. In order to prevent an increase in manufacturing costs, the Sb content may be 0.0001% or more, or may be 0.0010% or more.

(As: 0-0.0500%)
As is an element that is contained when scrap is used as a steel raw material, similar to Sn and Sb, and strongly segregates at grain boundaries, and the less the As, the better. In addition, excessive As content may lead to a decrease in bending strength during part deformation due to embrittlement of the steel sheet. The As content is 0.0500% or less, and may be 0.0400% or less. The As content may be 0%, but controlling the As content to less than 0.0001% may increase the refining time and increase the manufacturing cost. In order to prevent an increase in manufacturing costs, the As content may be 0.0001% or more, or may be 0.0010% or more.

(Mg: 0-0.0500%)
Mg is an element that can control the morphology of sulfides by adding a small amount, and is added as necessary. On the other hand, excessive Mg content forms coarse inclusions, which become the starting point for void generation, and may lead to a decrease in bending strength during part deformation. The Mg content is 0% or more, may be 0.0001% or more, may be 0.0010% or more, and may be 0.0500% or less, or may be 0.0400% or less.

(Ca: 0-0.0500%)
Ca is useful as a deoxidizing element and is also effective in controlling the morphology of sulfides. On the other hand, excessive Ca content may lead to a decrease in bending strength during part deformation due to embrittlement of the steel sheet. The Ca content is 0% or more, may be 0.0001% or more, may be 0.0010% or more, and may be 0.0500% or less, or may be 0.0400% or less.

(Y: 0-0.0500%)
Y is an element that can control the form of sulfides by adding a small amount, similar to Mg and Ca, and is added as necessary. On the other hand, excessive Y content generates coarse Y inclusions, which become the starting point of void generation, and may lead to a decrease in bending strength during part deformation. The Y content is 0% or more, may be 0.0001% or more, may be 0.0010% or more, and may be 0.0500% or less, or may be 0.0400% or less.

(Zr: 0-0.0500%)
Zr is an element that can control the form of sulfides by adding a small amount, similar to Mg, Ca, and Y, and is added as necessary. On the other hand, excessive Zr content generates coarse Zr inclusions, which become the starting point of void generation, and may lead to a decrease in bending strength during part deformation. The Zr content is 0% or more, may be 0.0001% or more, may be 0.0010% or more, and may be 0.0500% or less, or may be 0.0400% or less.

(La: 0-0.0500%)
La is an element that is effective in controlling the morphology of sulfides when added in small amounts, and is added as necessary. On the other hand, excessive La content generates La inclusions, which become the starting point for void generation, and may lead to a decrease in bending strength during part deformation. The La content is 0% or more, may be 0.0001% or more, may be 0.0010% or more, and may be 0.0500% or less, or may be 0.0400% or less.

(Ce: 0-0.0500%)
Ce is an element that can control the form of sulfides by adding a small amount, similar to La, and is added as necessary. On the other hand, excessive Ce content generates Ce inclusions, which become the starting point of void generation, and may lead to a decrease in bending strength during part deformation. The Ce content is 0% or more, may be 0.0001% or more, may be 0.0010% or more, and may be 0.0500% or less, or may be 0.0400% or less.

In the steel plate of this embodiment, the balance of the above-mentioned components is Fe and impurities. The impurities are components that are mixed in due to various factors in the manufacturing process, including raw materials such as ores and scraps, when the steel plate of this embodiment is industrially manufactured.

Next, the characteristics of the structure and properties of the steel plate according to an embodiment of the present invention will be described.

(Total area ratio of ferrite, pearlite and bainite: 0 to 30.0%)
Although ferrite, pearlite and bainite are effective structures for improving the strength-ductility balance of steel sheets, their inclusion in large amounts may lead to a decrease in local ductility. In addition, from the viewpoint of efficiently increasing the strength of steel, it is preferable that the area ratios of ferrite, pearlite and bainite are as small as possible. The total area ratio of ferrite, pearlite and bainite may be 0%, 1.0% or more, 30.0% or less, 25.0% or less, or 20.0% or less. Although the productivity is somewhat reduced, it is possible to make the total area ratio of ferrite, pearlite and bainite 0% by controlling the integrated manufacturing conditions with high precision.

(area ratio of retained austenite: 10.0 to 30.0%)
The retained austenite is an effective structure for improving the strength-ductility balance of a steel sheet. If the area ratio of the retained austenite is too small, the effect of increasing the strength due to the processing-induced transformation from the retained austenite to martensite cannot be obtained when the steel sheet is subjected to bending deformation, and this may result in a decrease in bending strength. On the other hand, if the area ratio of the retained austenite is too large, this may result in a decrease in bending strength as well as a decrease in yield strength. The area ratio of the retained austenite is 10.0% or more, and may be 13.0% or more, and may be 30.0% or less, or may be 25.0% or less.

In the steel plate according to the present embodiment, it is preferable that the steel structure of the steel plate contains needle-shaped retained austenite. The following effects can be expected by the retained austenite being "needle-shaped". That is, if the shape of the retained austenite is spherical (lumpy), deformation-induced transformation easily occurs with deformation of the steel plate, so bending deformation may begin at low stress. In contrast, if the shape of the retained austenite is needle-shaped, processing-induced transformation is less likely to occur, and bending strength is further increased. In the steel plate according to the present embodiment, the effect of the needle-shaped retained austenite and the effect of the difference in number density of precipitates are combined to significantly improve the bending strength of the steel plate. The area ratio of the needle-shaped retained austenite may be 30% or more or 50% or more, or 95% or less or 90% or less, when the area ratio of the entire retained austenite is 100%. In addition, in the present application, "needle-shaped retained austenite" refers to a ratio of the major axis to the minor axis (major axis/minor axis) of 3.0 or more. The "major axis" and "minor axis" of the retained austenite can be determined by observing the structure using EBSD. Specifically, in the observation of the structure, one retained austenite grain is identified, and the smallest Feret diameter of the grain is determined as the minor axis, and the largest Feret diameter of the grain is determined as the major axis.

(balance: fresh martensite and tempered martensite)
Fresh martensite and tempered martensite are extremely effective structures for increasing the strength of the steel sheet, and the higher the area ratio, the more preferable. In the steel sheet according to this embodiment, the remainder other than the above-mentioned ferrite, pearlite, bainite, and retained austenite is composed of fresh martensite and tempered martensite. The total area ratio of fresh martensite and tempered martensite may be 40.0% or more, 45.0% or more, 50.0% or more, 90.0% or less, or 85.0% or less. The area ratio of fresh martensite may be 5% or more, 10% or more, 20% or more, 30% or more, or 40% or more, and may be 80% or less, 70% or less, 60% or less, 50% or less, or 40% or less. Furthermore, the area ratio of tempered martensite may be 5% or more, 10% or more, 20% or more, 30% or more, or 40% or more, and may be 80% or less, 70% or less, 60% or less, 50% or less, or 40% or less.

(Difference in number density of precipitates in tempered martensite on the first surface on the front side of the steel sheet and the second surface on the back side: 0 to 10.0%)
The number density of precipitates in tempered martensite on the first surface on the front side and the second surface on the back side of the steel sheet is an important factor for increasing resistance to bending deformation. The higher the number density of both the first surface on the front side and the second surface on the back side of the steel sheet, the higher the bending strength. If the number density of one surface is high, strain is biased during bending deformation, and yielding occurs on the other surface with a low number density, resulting in a decrease in bending strength. Therefore, the smaller the difference between the number density of precipitates in tempered martensite on the first surface on the front side of the steel sheet and the number density of precipitates in tempered martensite on the second surface on the back side of the steel sheet, the more preferable it is. Specifically, it is important that the difference in number density of the precipitates is 10.0% or less. This difference in number density may be 8.0% or less, 6.0% or less, 4.0% or less, or 2.0% or less. In other words, in this embodiment, the ratio A1/A2 of the number density A1 of the precipitates on the first surface on the front side of the steel sheet to the number density A2 of the precipitates on the second surface on the back side of the steel sheet is 0.90 or more and 1.10 or less, and may be 0.92 or more, 0.94 or more, 0.96 or more, or 0.98 or more, or 1.08 or less, 1.06 or less, 1.04 or less, or 1.02 or less. In addition, making the difference in number density 0% is a factor that increases the manufacturing load and increases the manufacturing cost due to precise control of the steel sheet structure. In this respect, the difference in number density may be 0.1% or more. In addition, the precipitates are mainly carbides generated by tempering martensite, and the carbides may be cementite, or iron-based carbides or alloy carbides in which alloy elements such as Cr, Ti, and V are bonded to carbon instead of iron. The specific value of the number density of precipitates in the tempered martensite on the first surface on the front side and the second surface on the back side of the steel plate may be, for example, 1 precipitate/ ^μm2 or more, 5 precipitates/ ^μm2 or more, or 10 precipitates/μm2 ^{or more} , and may be 300 precipitates/μm2 or less, 100 precipitates/μm2 or ^less , or 30 precipitates/ ^μm2 or less.

In this application, for the sake of convenience, a distinction is made between the "front" and "back" of the steel plate, but there is no particular limitation as to which side of the steel plate is the front and which is the back.

(Yield strength YS)
In order to reduce the weight of a structure using steel as a material and to improve the strength at which plastic deformation begins, it is preferable that the yield strength of the steel material is high. On the other hand, if the yield strength is too high, the effect of shape change due to elastic deformation after plastic processing, so-called springback, may become large. The yield strength of the steel plate according to this embodiment may be 600 MPa or more, or 650 MPa or more. The upper limit of the yield strength is not particularly limited, but from the viewpoint of suppressing the effect of the above-mentioned springback, it may be 1100 MPa or less, or 1050 MPa or less.

(Tensile strength TS)
In order to reduce the weight of a structure using steel as a material and to improve the resistance of the structure to plastic deformation, it is preferable that the steel material has a large work hardening ability and exhibits maximum strength. On the other hand, if the tensile strength is too high, it may be prone to fracture with low energy during plastic deformation, and formability may be reduced. The tensile strength of the steel plate is not particularly limited, but may be 900 MPa or more, 980 MPa or more, and 2000 MPa or less, or 1800 MPa or less.

(Total elongation t-El)
When manufacturing a structure by cold forming a steel sheet as a raw material, elongation is required to finish the structure into a complex shape. If the total elongation is too low, the raw material may crack during cold forming. On the other hand, the higher the total elongation, the more preferable it is. However, if the total elongation is excessively increased, a large amount of retained austenite is required in the steel structure, which may reduce the yield strength during bending deformation. The total elongation of the steel sheet is not particularly limited, but may be 13% or more, 20% or more, 35% or less, or 30% or less.

(Hole expandability)
When manufacturing a structure by cold forming a steel sheet as a material, hole expandability is required as well as elongation in order to finish it into a complex shape. If the hole expandability is too small, the material may crack during cold forming. On the other hand, although the higher the hole expandability, the more preferable it is, if the hole expandability is excessively increased, a large amount of retained austenite is required in the steel structure, which may reduce the yield strength during bending deformation. The hole expansion ratio λ of the steel sheet is not particularly limited, but may be 20% or more, 25% or more, 90% or less, or 80% or less.

(Bendability)
When manufacturing a structure by cold forming a steel sheet as a material, bendability is also required to finish the structure into a complex shape. If the VDA bending angle is too small, the material may crack during cold forming. The higher the bendability, the better. The VDA bending angle of the steel sheet is not particularly limited, but may be 45° or more, or 50° or more.

(Thickness)
The plate thickness is a factor that affects the rigidity of the steel member after forming, and the greater the plate thickness, the higher the rigidity of the member. If the plate thickness is too small, the rigidity will decrease, and press formability may decrease due to the influence of unavoidable non-ferrous inclusions present inside the steel plate. On the other hand, if the plate thickness is too large, the press forming load will increase, leading to wear of the mold and reduced productivity. The plate thickness of the steel plate is not particularly limited, but may be 0.2 mm or more and 6.0 mm or less.

Next, we will describe the methods for observing and measuring the structure specified above, as well as the methods for measuring and evaluating the characteristics specified above.

(Method for measuring the total area ratio of ferrite, pearlite, and bainite)
The structure observation is performed with a scanning electron microscope. Prior to the observation, the sample for structure observation is polished by wet polishing with emery paper and diamond abrasive grains having an average particle size of 1 μm, and the observation surface is mirror-finished, and then the structure is etched with a 3% nitric acid alcohol solution. The magnification of the observation is 3000 times, and 10 random images of a 30 μm x 40 μm field of view at each 1/4 thickness position from the surface side of the steel plate are taken. The ratio of the structure is obtained by a point count method. A total of 100 lattice points arranged at intervals of 3 μm vertically and 4 μm horizontally are determined for the obtained structure image, the structure present under the lattice points is identified, and the structure ratio contained in the steel plate is obtained from the average value of the 10 sheets. Ferrite is a blocky crystal grain that does not contain iron-based carbides with a major axis of 100 nm or more inside. Bainite is a collection of lath-shaped crystal grains, and does not contain iron-based carbides with a major axis of 20 nm or more, or contains iron-based carbides with a major axis of 20 nm or more, and the carbides belong to a single variant, i.e., a group of iron-based carbides elongated in the same direction. Here, the group of iron-based carbides elongated in the same direction refers to iron-based carbides whose elongation directions differ by 5° or less. Bainite surrounded by grain boundaries with an orientation difference of 15° or more is counted as one bainite grain. Pearlite is a structure containing cementite precipitated in rows, and the area ratio is calculated by taking the area photographed with bright contrast in the secondary electron image as pearlite.

(How to distinguish between fresh martensite and tempered martensite)
Fresh martensite and tempered martensite are observed with scanning and transmission electron microscopes, and those containing Fe-based carbides (those with 1 Fe-based carbide/μm2 or ^more ) are identified as tempered martensite, and those containing almost no Fe-based carbides (those with less than 1 Fe-based carbide/ ^μm2 ) are identified as fresh martensite. Fe-based carbides having various crystal structures have been reported, but any Fe-based carbide may be contained. Depending on the heat treatment conditions, multiple types of Fe-based carbides may be present. In the present application, the total area ratio A1 of ferrite, pearlite, and bainite is measured by the above method, and the area ratio A2 of retained austenite is measured by the method described below, and the remainder obtained by subtracting the total value of the area ratios A1 and A2 from 100% is considered to be the total area ratio of fresh martensite and tempered martensite.

(Method of measuring area ratio of retained austenite)
The area fraction of retained austenite is determined by X-ray measurement as follows. First, a portion from the surface of the steel plate to 1/4 of the thickness of the steel plate is removed by mechanical polishing and chemical polishing, and the chemically polished surface is measured using MoKα rays as characteristic X-rays. Then, the area fraction of retained austenite in the center of the plate thickness is calculated using the following formula from the integrated intensity ratio of the diffraction peaks of (200) and (211) of the body-centered cubic lattice (bcc) phase and (200), (220) and (311) of the face-centered cubic lattice (fcc) phase.
Sγ=(I200f+I220f+I311f)/(I200b+I211b)×100
(Sγ is the area fraction of retained austenite in the center of the sheet thickness, I200f, I220f and I311f indicate the intensities of the diffraction peaks of (200), (220) and (311), respectively, of the fcc phase, and I200b and I211b indicate the intensities of the diffraction peaks of (200) and (211), respectively, of the bcc phase.)

To prepare a sample for X-ray diffraction, the steel plate is first reduced in thickness from the surface by mechanical polishing or the like to a specified thickness, and then strain is removed by chemical polishing or electrolytic polishing. At the same time, the sample is adjusted according to the above-mentioned method so that an appropriate surface is the measurement surface within the range of 1/8 to 3/8 of the plate thickness. Naturally, material anisotropy is further reduced by satisfying the above-mentioned X-ray intensity limitations not only in the vicinity of 1/4 of the plate thickness but for as many thicknesses as possible. However, by performing measurements within the range of 1/8 to 3/8 of the plate thickness from the surface, the material properties of the entire steel plate can be roughly represented. Therefore, the measurement range is set to 1/8 to 3/8 of the plate thickness.

Furthermore, the area ratio of acicular retained austenite can be measured, for example, by EBSD.

(Method for measuring number density of precipitates in tempered martensite on the first surface on the front side and the second surface on the back side of a steel sheet)
The number density of precipitates in tempered martensite on the first surface on the front side and the second surface on the back side of the steel sheet is measured as follows. First, a material for observation is taken from the front or back side of the steel sheet (meaning the front or back side of the base steel sheet. For example, in the case of a surface-treated steel sheet having a surface treatment layer such as plating, it means the front or back side of the base steel sheet excluding the surface treatment layer) at a depth position of 1/8 of the thickness in the sheet thickness direction, and prepared into a test piece for observation of a thin film or extraction replica. The test piece is observed at a magnification of 10,000 times with a transmission electron microscope to obtain images of at least 30 fields of view, and the number density of precipitates per unit area is measured in each observation image, and the value obtained by arithmetically averaging the number densities of the 30 fields of view is taken as the number density of precipitates on the first surface on the front side or the second surface on the back side. The field of view observed at a magnification of 10,000 times with a transmission electron microscope is a rectangular region with one side of about 600 nm, and the area of 30 fields of view for measuring the number density of precipitates is approximately 10.8 ^μm2 .

(Methods for measuring yield strength YP, tensile strength TS, and total elongation t-El)
The tensile test for measuring the yield strength, tensile strength and total elongation is performed in accordance with JIS Z 2241:2011, using a JIS No. 5 test piece taken from a direction in which the longitudinal direction of the test piece is parallel to the rolling direction perpendicular to the steel strip.

(Method of measuring hole expandability)
The hole expandability is evaluated by punching a circular hole having a diameter of 10 mm under the condition of a clearance of 12.5%, forming the hole with a 60 ° conical punch so that the burr is on the die side, and evaluating the hole expansion ratio λ (%). Five hole expansion tests are performed, and the average value is taken as the hole expansion ratio.

(Method of measuring bendability)
The bendability is evaluated by performing a test in accordance with the standard 238-100 of the German Association of the Automotive Industry (Verband der Automobilindustrie: VDA) using a test piece with a width of 60 mm, and the VDA bending angle is evaluated by measuring the maximum bending angle α. The bending strength is evaluated as the load at a bending angle of 5° divided by the plate thickness.

<Method of manufacturing steel sheet>
The method for producing a steel sheet according to this embodiment is characterized by using a material having the above-mentioned chemical composition and consistently managing hot rolling, cold rolling, and annealing. Specifically, the method for producing a steel sheet according to this embodiment is characterized by including the steps of hot rolling a steel billet (steel slab) having the same chemical composition as that described above for the steel sheet, coiling it, pickling the obtained hot-rolled sheet, cold rolling it, annealing it, performing aging treatment, and then annealing it again. More specifically, the method for producing a steel sheet according to this embodiment includes the steps of:
hot rolling the steel slab having the above chemical composition to obtain a hot rolled sheet;
coiling the hot-rolled sheet;
pickling the hot-rolled sheet;
cold rolling the hot rolled sheet to obtain a cold rolled sheet;
The cold-rolled sheet is subjected to a first annealing (Q annealing);
The cold-rolled sheet subjected to the first annealing is subjected to a second annealing (IA annealing); and
performing an aging treatment between the first annealing and the second annealing;
Including,
The first annealing is a process of heating the cold-rolled sheet to a temperature in the austenite single phase region and 1000° C. or less, and then cooling the sheet to obtain a martensite structure having an area ratio of 90.0% or more.
The second annealing is a process of holding the cold-rolled sheet in a two-phase region of ferrite and austenite to obtain retained austenite,
The aging treatment is characterized by including an aging treatment 1 in which one of the front side and the back side of the cold-rolled sheet is subjected to a tensile deformation with a bending R of 2.0 m or less and held at 0 to 40° C. for 20 hours or more, and an aging treatment 2 in which the other of the front side and the back side of the cold-rolled sheet is subjected to a tensile deformation with a bending R of 2.0 m or less and held at 0 to 40° C. for 20 hours or more. Each step will be described in detail below, focusing on the key points of this embodiment.

(Hot rolling finishing temperature)
In this embodiment, a hot-rolled sheet is obtained by hot rolling a steel slab obtained by a known method such as a continuous casting method. Here, the finish rolling temperature of the hot rolling is a factor that has an effect on controlling the texture of the prior austenite grain size. From the viewpoint that the rolling texture of austenite develops and causes anisotropy in the steel material properties, the finish rolling temperature is preferably 650°C or higher, and from the aim of suppressing the deviation of the texture due to abnormal grain growth of austenite, the finish rolling temperature is preferably 950°C or lower.

(Coil winding temperature)
The temperature at which the hot-rolled sheet is coiled (the coiling temperature of the hot-rolled coil) is a factor that controls the state of oxide scale formation in the hot-rolled sheet and affects the strength of the hot-rolled sheet. The thinner the thickness of the scale formed on the surface of the hot-rolled sheet, the better, and therefore the lower the coiling temperature is. In addition, when the coiling temperature is extremely lowered, special equipment is required. In addition, if the coiling temperature is too high, as described above, the oxide scale formed on the surface of the hot-rolled sheet becomes significantly thick. From the above viewpoints, the temperature at which the hot-rolled sheet is coiled may be 700° C. or less, 680° C. or less, 0° C. or more, or 20° C. or more.

(Pickling of hot-rolled sheet)
The pickling of the hot-rolled sheet is aimed at removing scales and the like, and may be carried out under known pickling conditions.

(Reduction in cold rolling)
In cold rolling, if the total reduction is too large, the ductility of the base steel sheet is lost, and the risk of the base steel sheet breaking during cold rolling increases. In this respect, the total reduction in cold rolling is preferably 85% or less. On the other hand, in order to sufficiently promote recrystallization in the annealing process, the total reduction is preferably 20% or more, and more preferably 30% or more. Annealing may be performed at a temperature of 700°C or less in order to reduce the cold rolling load before cold rolling.

(Holding temperature in first annealing)
In the first annealing (Q annealing), the cold-rolled sheet, which is the base steel sheet, is heated to a temperature of Ac3 point or higher and 1000°C or lower (i.e., austenite single phase region and 1000°C or lower). The reason for setting the maximum heating temperature to Ac3 point or higher is that by heating the base steel sheet to the austenite single phase region, a martensite structure with an area ratio of 90% or more can be obtained by subsequent quenching, and the precipitation of ε carbide by aging can be promoted. If the steel sheet is held at a lower temperature than this, a martensite-based structure cannot be obtained, and the bending strength is significantly reduced. On the other hand, if the steel sheet is heated to a temperature of more than 1000°C, the surface layer of the steel sheet is decarburized, and the strength is reduced, so that the bending strength may be reduced.

(Holding time in first annealing)
In the first annealing (Q annealing), it is preferable to hold the steel sheet at a heating temperature of Ac3 point or higher and 1000°C or lower for 5 seconds or more. If the holding time is too short, the austenite transformation of the base steel sheet will not proceed sufficiently, and in addition, the concentration of substitutional elements such as Mn that stabilize austenite will not be sufficient in the austenite, so that the retained austenite will become unstable and the ductility of the steel sheet may decrease significantly. From these viewpoints, the holding time is more preferably 10 seconds or more. More preferably, it is 20 seconds or more.

(Atmosphere in the first annealing)
In the first annealing (Q annealing), the oxygen potential in one or both of the heating zone and the soaking zone during annealing may be controlled in order to provide a decarburized layer in the surface layer of the steel sheet to improve bendability. Specifically, the annealing is preferably performed in an atmosphere containing 0.1 to 30% by volume of hydrogen and H 2 O with a dew point of −40 to 20° C., with the balance being nitrogen and impurities. More preferably, the atmosphere contains 0.5 to 20% by volume of hydrogen and _{H 2} O with a dew point of −30 to 15° C., and even more preferably, the atmosphere contains 1 to 10% by volume of hydrogen and H ₂ O with a dew point of −20 to 10° C.

(Cooling rate in first annealing)
In the first annealing (Q annealing), when cooling after heating and holding, it is preferable to cool from 750 ° C. to 550 ° C. at an average cooling rate of 100 ° C. / s or less. The lower limit of the average cooling rate is not particularly limited as long as a martensite structure of 90% or more in area ratio is obtained, but may be, for example, 3 ° C. / s. The reason for setting the lower limit of the average cooling rate to 3 ° C. / s is to suppress the occurrence of ferrite transformation in the base steel sheet and the area ratio of martensite in the steel structure after Q annealing to be less than 90%. More preferably, it is 10 ° C. / s or more, more preferably 15 ° C. / s or more, and even more preferably 20 ° C. / s or more. On the other hand, if the cooling rate from 750 ° C. to 550 ° C. is too fast, a low-temperature transformation structure will also occur in the steel sheet surface layer, causing hardness variation. In this regard, the average cooling rate is preferably 100 ° C. / s or less, more preferably 80 ° C. / s or less, and even more preferably 50 ° C. / s or less. The cooling rate is not limited at temperatures above 750° C. because ferrite transformation is unlikely to occur at these temperatures. Also, the cooling rate is not limited at temperatures below 550° C. because a low-temperature transformed structure is obtained at these temperatures.

(Cooling stop temperature and reheating in the first annealing)
After the above cooling, the material may be further cooled to a temperature of 25° C. to 550° C., and then allowed to dwell in a temperature range of 150° C. to 550° C. The reason why the lower limit of the cooling stop temperature is set to 25° C. is that excessive cooling not only requires a significant investment in equipment, but also has a saturating effect. The dwell time is not particularly limited, and may be, for example, 30 seconds to 500 seconds.

(Aging Treatment 1)
Holding the steel sheet, which has been controlled to have a martensite-based structure by the first annealing, at 0 to 40° C. for 20 hours or more while bending the steel sheet with a bending radius R of 2.0 m or less is an important factor in increasing the bending strength of the steel sheet. During this treatment, the carbon atoms dissolved in the martensite form clusters or transition carbides, which become the nuclei of carbide precipitation when the temperature is raised in the subsequent second annealing. In order to finely disperse carbides and increase the bending strength, it is important to have fine and dense clusters or transition carbides that become the nuclei of carbide precipitation. In order to promote the generation of clusters or transition carbides, the use of tensile strain is extremely effective, and this effect is easily obtained in bending deformation with a bending radius R of 2.0 m or less. The bending radius R may be 1.8 m or less, 1.5 m or less, or 1.3 m or less. On the other hand, if the bending radius R exceeds 2.0 m, it becomes difficult to obtain this effect. For example, the steel sheet (steel strip) after the first annealing is wound into a coil, whereby the above bending deformation can be imparted to the steel sheet.

In addition, when the holding temperature is less than 0°C, the clustering of carbon atoms or the formation of transition carbides is suppressed, and when the holding temperature exceeds 40°C, the transition carbides are formed coarsely (the number of nuclei decreases), so that it is difficult to obtain fine carbides in the second annealing, and bending strength may decrease. If the holding temperature in aging treatment 1 and aging treatment 2 described below is within the range of 0 to 40°C, the difference in the number of precipitates between aging treatment 1 and aging treatment 2 becomes small, and the difference between the number density of precipitates on the first surface on the front side of the steel sheet and the number density of precipitates on the second surface on the back side of the steel sheet is within 10%. The holding temperature may be 5°C or more, 10°C or more, 35°C or less, or 30°C or less.

Furthermore, if the holding time is less than 20 hours, the number of nuclei generated is not stable, and a sufficient amount of nuclei is not generated, and it may be difficult to suppress the difference between the number density of precipitates on the first surface on the front side of the steel sheet and the number density of precipitates on the second surface on the back side of the steel sheet to within 10%. A longer holding time is preferable, and may be 30 hours or more, 40 hours or more, or 50 hours or more. Note that if the holding time exceeds 300 hours, the clustering of carbon atoms or the formation of transition carbides will be saturated, and if the holding time is longer than this, it will be difficult for a significant change to occur in the form (size) of the precipitates, so the holding time may be 300 hours or less. Note that, although the precipitates will become larger with a longer holding time, there will be no significant change in the number of precipitates. That is, when the holding time of aging treatment 1 and aging treatment 2 described below is 20 hours or longer, the number of nuclei generated is stabilized, and the difference between the number density of precipitates on the first surface on the front side of the steel sheet and the number density of precipitates on the second surface on the back side of the steel sheet is within 10%.

(Aging Treatment 2)
When bending deformation is applied to a steel sheet to promote aging, clustering of carbon atoms and precipitation of transition carbides occur significantly in the region subjected to tensile deformation, so that if the winding and unwinding process is performed only once by the above-mentioned aging treatment 1, precipitates will be finely dispersed only on one side of the steel sheet. For this reason, in the manufacturing method according to this embodiment, after the aging treatment 1 in which precipitates are precipitated and dispersed on one side of the front side and the back side of the steel sheet, aging treatment 2 in which precipitates are precipitated and dispersed on the other side of the front side and the back side of the steel sheet is performed. For example, after the first annealing, the sheet is wound into a coil so that the front side of the sheet is the outside and the back side is the inside, and tensile deformation with a bending radius R of 2.0 m or less is applied to the front side of the sheet to perform aging treatment 1, and then the coil is unwound, and the sheet is again wound into a coil so that the back side of the sheet is the outside and the front side is the inside, and tensile deformation with a bending radius R of 2.0 m or less is applied to the back side of the sheet to perform aging treatment 2. In aging treatment 2, similarly to the holding conditions in aging treatment 1, by performing bending deformation with a bending radius R of 2.0 m or less, the difference in number density of precipitates on the first surface on the front side of the steel sheet and the difference in number density of precipitates on the second surface on the back side of the steel sheet can be suppressed to within 10%. The bending radius R may be 1.8 m or less, 1.5 m or less, or 1.3 m or less. On the other hand, if the bending radius R exceeds 2 m, it becomes difficult to suppress the difference in number density of precipitates to within 10%.

In addition, in the aging treatment 2, as in the aging treatment 1, when the holding temperature is less than 0°C, the clustering of carbon atoms or the formation of transition carbides is suppressed, and when the holding temperature exceeds 40°C, the transition carbides are formed coarsely (the number of nuclei decreases), so that it is difficult to obtain fine carbides in the second annealing, and bending strength may decrease. As described above, if the holding temperature in the aging treatment 1 and the aging treatment 2 described later is within the range of 0 to 40°C, the difference between the number density of precipitates on the first surface on the front side of the steel sheet and the number density of precipitates on the second surface on the back side of the steel sheet is within 10%. The holding temperature may be 5°C or more, 10°C or more, 35°C or less, or 30°C or less.

Furthermore, in the aging treatment 2, as in the aging treatment 1, if the holding time is less than 20 hours, the number of nuclei generated is not stable, and it may be difficult to suppress the difference between the number density of precipitates on the first surface on the front side of the steel sheet and the number density of precipitates on the second surface on the back side of the steel sheet to within 10%. A longer holding time is preferable, and may be 30 hours or more, 40 hours or more, or 50 hours or more. As in the aging treatment 1, if the holding time exceeds 300 hours, the clustering of carbon atoms or the formation of transition carbides becomes saturated, and if the holding time is longer than that, it is difficult for a significant change in the morphology (size) of the precipitates to occur, so the holding time may be 300 hours or less.

(Holding temperature in second annealing)
In the second annealing (IA annealing), the holding temperature is a temperature at which the two-phase region of ferrite and austenite is formed. For example, it is preferably 720°C or more and 860°C or less. If the annealing temperature is less than 720°C, austenite is not sufficiently generated. In this case, the martensite obtained in the first annealing (Q annealing) is tempered, which leads to the precipitation of carbides, and the area ratio of the predetermined retained austenite may not be satisfied. In addition, since the area ratio of austenite at the maximum heating temperature (annealing temperature) also decreases, it becomes impossible to concentrate the carbon necessary to obtain the retained austenite in the austenite, and it may become impossible to secure 10.0% or more of the retained austenite. On the other hand, if the annealing temperature exceeds 860°C, austenite is excessively generated, and the bending strength decreases due to the reduction of tempered martensite containing precipitates, and it may become impossible to secure 10.0% or more of the retained austenite. Therefore, it is preferable that the upper limit of the holding temperature in the second annealing is 860°C. The annealing may be carried out in air, or in order to improve the adhesion of the plating, it may be carried out in an atmosphere in which the hydrogen concentration and dew point are controlled.

(Holding time in second annealing)
In the second annealing (IA annealing), it is preferable to hold the steel sheet at a heating temperature of 720°C or more and 860°C or less for 5 seconds or more. If the holding time is too short, the austenite transformation of the base steel sheet will not proceed sufficiently, and in addition, the concentration of substitutional elements such as Mn that stabilize austenite will not be sufficient in the austenite, so that the retained austenite will become unstable and the ductility of the steel sheet may decrease significantly. From these viewpoints, the holding time is more preferably 10 seconds or more. More preferably, it is 20 seconds or more.

(Atmosphere in second annealing)
In the second annealing (IA annealing), similarly to the first annealing (Q annealing), the oxygen potential in one or both of the heating zone and the soaking zone during annealing may be controlled in order to provide a decarburized layer in the surface layer of the steel sheet to improve the bendability. Specifically, the annealing is preferably performed in an atmosphere containing 0.1 to 30% by volume of hydrogen and H ₂ O with a dew point of −40 to 20° C., with the balance being nitrogen and impurities. More preferably, the atmosphere contains 0.5 to 20% by volume of hydrogen and H ₂ O with a dew point of −30 to 15° C., and even more preferably, the atmosphere contains 1 to 10% by volume of hydrogen and H ₂ O with a dew point of −20 to 10° C.

(Cooling rate in second annealing)
In the second annealing (IA annealing), when cooling after heating and holding, it is preferable to cool from 750 ° C. to 550 ° C. at an average cooling rate of 100 ° C. / s or less. The lower limit of the average cooling rate is not particularly limited, but may be, for example, 2.5 ° C. / s. The reason for setting the lower limit of the average cooling rate to 2.5 ° C. / s is to suppress the softening of the base steel sheet by the occurrence of ferrite transformation from acicular austenite in which alloy elements are concentrated in the base steel sheet. If the average cooling rate is too slow, the strength is likely to decrease. More preferably, it is 5 ° C. / s or more, more preferably 10 ° C. / s or more, and even more preferably 20 ° C. / s or more. On the other hand, if the cooling rate from 750 ° C. to 550 ° C. is too fast, a low-temperature transformation structure is also generated in the surface layer of the steel sheet, which causes hardness variation. In this respect, the average cooling rate is preferably 100 ° C. / s or less, more preferably 80 ° C. / s or less, and even more preferably 50 ° C. / s or less. The cooling rate is not limited at temperatures above 750° C. because ferrite transformation is unlikely to occur at these temperatures. Also, the cooling rate is not limited at temperatures below 550° C. because a low-temperature transformed structure is obtained at these temperatures.

(Cooling stop temperature and reheating in the second annealing)
After the above cooling, the steel sheet may be further cooled to a temperature of 25°C to 550°C, and then reheated and retained in a temperature range of 150°C to 550°C. When the steel sheet is cooled in the above temperature range, martensite is generated from untransformed austenite during cooling. Then, by reheating, carbon is concentrated from martensite to untransformed austenite, improving the strength-ductility balance of the steel sheet. The reason why the lower limit of the cooling stop temperature is set to 25°C is that excessive cooling not only requires a significant capital investment, but also saturates its effect. The retention time is not particularly limited, but may be, for example, 30 seconds to 500 seconds.

(Conditions for obtaining acicular retained austenite in the second annealing)
In the method for producing a steel sheet according to the present embodiment, it is preferable to obtain needle-shaped retained austenite in the second annealing (IA annealing). For example, by controlling the temperature so that the area ratio of the retained austenite obtained through IA annealing is 10 to 50% and the temperature change of the steel sheet per second in the holding step of IA annealing is within ±3°C, alloy elements are segregated at the interface between ferrite and austenite during holding in the two-phase region to reduce the mobility of the interface, and needle-shaped retained austenite can be obtained even at room temperature.

(Residence temperature)
Furthermore, after reheating and before immersion in the plating bath, the steel sheet may be retained in a temperature range of 350 to 550°C. Retention in this temperature range not only contributes to tempering of martensite, but also eliminates temperature unevenness in the width direction of the sheet, improving the appearance after plating. If the cooling stop temperature is 350 to 550°C, retention may be performed without reheating.

(Residence Time)
The retention time is preferably 30 seconds or more and 300 seconds or less in order to obtain the desired effect.

(Tempering)
In a series of annealing steps, after cooling the cold-rolled sheet or the plated steel sheet obtained by plating the cold-rolled sheet to room temperature, or during cooling to room temperature (but below Ms), reheating may be started, and the temperature may be held in a temperature range of 150°C to 400°C for 2 seconds or more. According to this process, the martensite generated during cooling after reheating is tempered to form tempered martensite, thereby improving hydrogen embrittlement resistance. When the tempering process is performed, if the holding temperature is too low or the holding time is too short, the martensite is not sufficiently tempered, and there is almost no change in the microstructure and mechanical properties. On the other hand, if the holding temperature is too high, the dislocation density in the tempered martensite decreases, leading to a decrease in tensile strength. Therefore, when tempering is performed, it is preferable to hold the temperature range of 150°C to 400°C for 2 seconds or more. Tempering may be performed in a continuous annealing facility, or may be performed offline after continuous annealing in a separate facility. In this case, the tempering time varies depending on the tempering temperature. That is, the lower the temperature, the longer the time, and the higher the temperature, the shorter the time.

(Plating)
The steel sheet may be heated or cooled to (galvanizing bath temperature -40) ° C. to (galvanizing bath temperature +50) ° C. as necessary to perform hot-dip galvanization. A hot-dip galvanized layer is formed on the surface of the steel sheet by the hot-dip galvanization process. In this case, the corrosion resistance of the cold-rolled sheet is improved, which is preferable. In this embodiment, the type of the plating layer is not limited to a hot-dip galvanized layer, and various coating layers can be adopted. In addition, the timing of plating the surface of the steel sheet is not particularly limited. For example, in the manufacturing method according to this embodiment, in the IA annealing, a coating layer made of zinc, aluminum, magnesium, or an alloy thereof may be formed on the front and back surfaces of the steel sheet in the process of cooling to room temperature after holding the cold-rolled sheet in the two-phase region of ferrite and austenite. Alternatively, the coating layer may be formed on the front and back surfaces of the steel sheet after annealing.

(Steel sheet temperature when immersed in plating bath)
The temperature of the steel sheet when immersed in the hot-dip galvanizing bath is preferably in the range from a temperature 40°C lower than the hot-dip galvanizing bath temperature (hot-dip galvanizing bath temperature -40°C) to a temperature 50°C higher than the hot-dip galvanizing bath temperature (hot-dip galvanizing bath temperature +50°C). If this temperature is below the hot-dip galvanizing bath temperature -40°C, heat removal during immersion in the galvanizing bath is large, and part of the molten zinc may solidify, deteriorating the appearance of the coating. If the sheet temperature before immersion is below the hot-dip galvanizing bath temperature -40°C, further heating may be performed by any method before immersion in the galvanizing bath, and the sheet temperature may be controlled to the hot-dip galvanizing bath temperature -40°C or higher before immersion in the galvanizing bath. In addition, if the temperature of the steel sheet when immersed in the galvanizing bath exceeds the hot-dip galvanizing bath temperature +50°C, operational problems due to an increase in the galvanizing bath temperature may be induced.

(Composition of plating bath)
The composition of the plating bath is preferably Zn-based, with an effective Al content (the total Al content minus the total Fe content in the plating bath) of 0.050 to 0.250 mass%. If the effective Al content in the plating bath is too small, Fe may penetrate excessively into the plating layer, resulting in a decrease in plating adhesion. On the other hand, if the effective Al content in the plating bath is too large, Al-based oxides that inhibit the movement of Fe and Zn atoms may be generated at the boundary between the steel sheet and the plating layer, resulting in a decrease in plating adhesion. The effective Al content in the plating bath is more preferably 0.065 mass% or more, and more preferably 0.180 mass% or less.

(Steel sheet temperature after immersion in plating bath)
When alloying treatment is performed on the hot-dip galvanized layer, the steel sheet on which the hot-dip galvanized layer is formed is heated to a temperature range of 450 to 600°C. If the alloying temperature is too low, the alloying may not proceed sufficiently. On the other hand, if the alloying temperature is too high, the alloying may proceed too much, and the Fe concentration in the plating layer may exceed 15% due to the formation of Γ phase, which may deteriorate the corrosion resistance. The alloying temperature is more preferably 470°C or higher, and more preferably 550°C or lower. The alloying temperature needs to be changed depending on the component composition of the steel sheet, so it may be set while checking the Fe concentration in the plating layer.

(Pre-processing)
In order to further improve the plating adhesion, the steel sheet may be plated with one or more of Ni, Cu, Co and Fe before annealing in a continuous hot-dip galvanizing line or the like.

(Post-processing)
For the purpose of improving paintability and weldability, a top layer plating may be applied to the plating surface of a hot-dip galvanized steel sheet having a steel sheet and a hot-dip galvanizing coating formed on the surface of the steel sheet, or a plating surface of a galvannealed steel sheet having a steel sheet and a hot-dip galvanizing coating formed on the surface of the steel sheet, or various treatments such as chromate treatment, phosphate treatment, lubricity improvement treatment, weldability improvement treatment, etc.

(Skin pass rolling ratio)
Furthermore, skin pass rolling may be performed for the purpose of correcting the steel sheet shape and introducing mobile dislocations to improve ductility. The reduction ratio of the skin pass rolling after heat treatment is preferably in the range of 0.1 to 1.5%. If it is less than 0.1%, the effect is small and control is difficult, so this is the lower limit. If it exceeds 1.5%, productivity drops significantly, so this is the upper limit. Skin pass rolling may be performed inline or offline.

<Additional Information>
The difference in hardness between the front and back of the steel sheet is not substantially related to the difference in number density of precipitates between the front and back of the steel sheet. That is, even if the difference in hardness between the front and back of the steel sheet is reduced, the difference in number density of precipitates between the front and back of the steel sheet cannot be reduced, the bending strength of the steel sheet is not necessarily improved, and the bending strength difference between the front and back of the steel sheet cannot be reduced. By reducing the difference in number density of precipitates between the front and back of the steel sheet as in the steel sheet according to this embodiment, the bending strength of the steel sheet can be improved and the bending strength difference between the front and back of the steel sheet can be reduced. In order to reduce the difference in number density of precipitates between the front and back of the steel sheet, it is effective to perform aging treatments 1 and 2 between the first annealing (Q annealing) and the second annealing (IA annealing) for the cold-rolled sheet as in the manufacturing method of the steel sheet according to this embodiment. Conventionally, there has been no consideration on reducing the difference in number density of precipitates between the front and back of the steel sheet, and no consideration was given to performing Q annealing, aging treatments 1 and 2, and IA annealing as in the manufacturing method according to this embodiment.

The following is an example of the present invention. The present invention is not limited to this example of conditions. The present invention allows for various conditions to be adopted as long as the gist of the invention is not exceeded and the object is achieved.

(Example 1)
Steels having various chemical compositions were melted to produce steel billets. These billets were inserted into a furnace heated to 1220°C, and were subjected to a homogenization treatment of holding for 60 minutes, after which they were taken out into the air and hot rolled to obtain steel plates with a thickness of 2.8 mm. In the hot rolling, the finishing rolling was completed at a temperature of 910°C, and the plates were cooled to 550°C and wound up. Subsequently, the oxide scale of the hot rolled plates was removed by pickling, and cold rolling was performed at a rolling reduction of 45.0% to finish the plates to a thickness of 1.54 mm. Furthermore, the cold rolled plates were Q-annealed, specifically, the temperature was raised to 930°C, and the holding time in that temperature range was set to 90 seconds. Next, the cold rolled plates were cooled and held at 280°C, and wound up into a coil shape with a maximum radius of 1.4 m. The area ratio of martensite in the steel plates after winding was 90% or more for all steel components. The coil after winding was subjected to aging treatment 1 in which the coil was held at a temperature range of 6°C to 22°C for 38 hours, and then the coil was unwound and wound up to form a coil with a maximum radius of 1.4 m again. The steel sheet was bent in the opposite direction to that of aging treatment 1, and was again held at a temperature range of 6°C to 22°C for 38 hours. The aged steel sheet subjected to the two aging treatments was then subjected to IA annealing, specifically, the temperature was raised to 785°C, and the holding time in that temperature range was set to 130 seconds. Next, the aged steel sheet was cooled to 270°C, reheated to 390°C, and held at that temperature for 140 seconds, and then cooled to room temperature, after which skin pass rolling was performed. The chemical compositions of the samples obtained by analyzing the samples taken from each steel sheet are as shown in Tables 1 to 3. In Tables 1 to 3, "-" means that the content was below the detection limit. The balance other than the components shown in Tables 1 to 3 is Fe and impurities. Table 4 shows the evaluation results of the properties of the steel sheets that were subjected to the above-mentioned thermomechanical treatment.

In Table 4, the methods for measuring the "area ratio," "yield strength YS," "tensile strength TS," "total elongation t-El," "hole expansion ratio λ," and "difference in number density of precipitates in tempered martensite" of each structure/phase in the steel structure are as described above. For "yield strength YS," a value of 600 MPa or more was judged to be "pass." For "bending strength (bending resistance)," two evaluation indices were used: (1) the bending strength value of the steel plate itself, and (2) the difference in bending strength between the front and back surfaces of the steel plate. Of these, the value of (1) "bending strength" was judged based on the load during VDA bending described above. Specifically, bending tests were performed on both the front and back sides of the steel plate, and the load when a 5° bending angle was applied was rated as "A" if it was 1400 N or more per 1 mm of plate thickness on both sides, "B" if it was 900 N or more but less than 1400 N, and "C" if it was less than 900 N. Those that were A or B were judged as "passed". Furthermore, (2) "Difference in bending strength between the front and back sides of the steel plate" was also judged based on the load during VDA bending. Specifically, those that had a difference in load between the front and back sides when a predetermined bending was applied to each of the front and back sides of the steel plate were judged as "A" if it was within 3%, "B" if it was more than 3% but not more than 8%, and "C" if it was more than 8%, and those that were A or B were judged as "passed".

The results shown in Tables 1 to 4 reveal the following:

Because the C content in the steel plate of BA-1 was too low, the yield strength YS and tensile strength TS of the steel were reduced, and sufficient bending resistance (bending yield strength) could not be ensured.

BB-1 had too much C content in the steel plate, which reduced elongation and led to brittle fracture of the steel, resulting in reduced bending resistance.

It is believed that in BC-1, the Si content in the steel plate was too low, which reduced the stability of the retained austenite structure against processing and made it impossible to suppress the precipitation of carbides in the martensite structure during aging. As a result, the total area ratio of ferrite, pearlite, and bainite increased, reducing the yield strength YS and tensile strength TS of the steel plate, and the area ratio of retained austenite decreased, making it impossible to obtain the effect of increasing strength due to processing-induced transformation from retained austenite to martensite when bending deformation is applied to the steel plate, and also making it impossible to ensure sufficient bending resistance.

It is believed that the formation of ε carbides was suppressed during aging treatment because the Si content in the steel plate of BD-1 was too high. As a result, sufficient bending resistance could not be ensured.

Because the Mn content in the steel plate of BE-1 was too low, ferrite transformation was likely to occur during the cooling process of Q-annealing, the martensite structure ratio after Q-annealing decreased, and the area ratio of retained austenite in the final steel plate also decreased. As a result, the yield strength YS and tensile strength TS of the steel plate decreased, and the effect of increasing strength due to the processing-induced transformation from retained austenite to martensite when bending deformation was applied to the steel plate could not be obtained, and sufficient bending resistance could not be ensured.

In BF-1, the Mn content in the steel plate was too high, causing a noticeable Mn-enriched layer due to micro-segregation and central segregation in the steel plate, and because differences in the solidification rates on the front and back sides of the slab caused differences in the distribution of the Mn-enriched layer on the front and back sides of the steel plate, differences in the formation of Mn segregation bands also caused differences in the number density of precipitates in the tempered martensite, resulting in differences in the bending strength on the front and back sides. As a result, sufficient bending resistance could not be ensured.

It is believed that the P content in the steel plate of BG-1 was too high, which led to brittle fracture of the steel plate and promoted a decrease in bending strength during bending deformation. As a result, sufficient bending resistance could not be secured.

The S content in the steel plate of BH-1 was too high, which is thought to have led to the formation of non-metallic inclusions, reducing the ductility of the steel plate and causing voids to form from the non-metallic inclusions during bending deformation. As a result, sufficient bending resistance could not be ensured.

In BI-1, the Al content in the steel plate was too high, which is thought to have excessively promoted the ferrite transformation and bainite transformation during the cooling process during annealing. As a result, the yield strength YS and tensile strength TS of the steel plate decreased, and sufficient bending resistance could not be ensured.

The steel plate BJ-1 had too much N, which combined with Al to form a large amount of AlN, and it is believed that these nitrides became the starting points for the generation of voids during bending deformation. As a result, sufficient bending resistance could not be ensured.

The Ti content in the steel plate of BK-1 was too high, which resulted in the precipitation of large coarse carbides, and it is believed that these carbides became the starting points for the generation of voids during bending deformation. As a result, sufficient bending resistance could not be ensured.

Because the Co content in the steel plate of BL-1 was too high, many fine Co carbides precipitated, and it is believed that these carbides became the starting points for the generation of voids during bending deformation. As a result, sufficient bending resistance could not be ensured.

Because the Ni content in the steel plate of BM-1 was too high, it is believed that this affected the peeling properties of the oxide scale during hot rolling and promoted the occurrence of scratches on the steel plate surface. As a result, sufficient bending resistance could not be ensured.

The Mo content in the steel plate of BN-1 was too high, causing the precipitation of many fine Mo carbides, which are believed to have become the starting points for the generation of voids during bending deformation. As a result, sufficient bending resistance could not be ensured.

In BO-1, the Cr content in the steel plate was too high, which promoted the formation of retained austenite, and the presence of excessive retained austenite meant that sufficient bending resistance could not be ensured.

Because the oxygen content in the steel plate of BP-1 was too high, a large amount of oxides were generated as inclusions, and these oxides on the punched and cut edges formed notch-like scratches and large dimples on the edges, which are thought to have led to stress concentration during bending deformation and become the starting point for crack formation. As a result, sufficient bending resistance could not be ensured.

The B content in the steel plate of BQ-1 was too high, which led to the formation of coarse B inclusions in the steel, and it is believed that these inclusions became the starting points for the generation of voids. As a result, sufficient bending resistance could not be ensured.

The Nb content in the steel plate of BR-1 was too high, causing the precipitation of many fine, hard Nb carbides, which are believed to have become the starting points for the generation of voids. As a result, sufficient bending resistance could not be ensured.

The V content in the BS-1 steel plate was too high, which led to the precipitation of a lot of carbonitrides, and it is believed that these carbonitrides became the starting point for the generation of voids. As a result, sufficient bending resistance could not be ensured.

It is believed that BT-1 had too much Cu in the steel plate, which increased the strength of the steel plate while decreasing its ductility. As a result, the bending strength during bending deformation decreased, and sufficient bending resistance could not be secured.

The W content in the steel plate of BU-1 was too high, which is thought to have facilitated the development of voids starting from coarse carbides. As a result, sufficient bending resistance could not be ensured.

In BV-1, the Ta content in the steel plate was too high, which is thought to have caused a large number of fine Ta carbides to precipitate, making it easier for voids to form from these carbides. As a result, sufficient bending resistance could not be ensured.

In BW-1, the Sn content in the steel plate was too high, which is thought to have caused the steel plate to become embrittled, reducing its bending strength during bending deformation, and sufficient bending resistance could not be ensured.

In BX-1, the Sb content in the steel plate was too high, which is thought to have caused the steel plate to become embrittled, reducing the bending strength during bending deformation, and sufficient bending resistance could not be ensured.

Because the As content in the steel plate of BY-1 was too high, it is believed that the embrittlement of the steel plate reduced its bending strength during bending deformation, and sufficient bending resistance could not be ensured.

The Mg content in the steel plate of BZ-1 was too high, which resulted in the formation of coarse inclusions, which are believed to have become the starting points for the generation of voids. As a result, sufficient bending resistance could not be ensured.

CA-1 had too much Ca content in the steel plate, which is thought to have caused the steel plate to become embrittled, reducing its bending strength during bending deformation, and it was not possible to ensure sufficient bending resistance.

The Y content in the steel plate of CB-1 was too high, which resulted in the formation of large Y inclusions, which are believed to have become the starting points for the generation of voids. As a result, sufficient bending resistance could not be secured.

In CC-1, the Zr content in the steel plate was too high, which resulted in the formation of coarse Zr inclusions, which are believed to have become the starting points for the generation of voids. As a result, sufficient bending resistance could not be secured.

CD-1 steel plate had too much La content, which is thought to have caused La inclusions to form and these inclusions became the starting points for the generation of voids. As a result, sufficient bending resistance could not be secured.

In CE-1, the Ce content in the steel plate was too high, which is thought to have caused the formation of Ce inclusions, which became the starting point for the generation of voids. As a result, sufficient bending resistance could not be ensured.

In contrast, for A-1 to AZ-1, the steel sheets having the specified chemical composition were produced under the specified conditions, and thus the specified metal structure was obtained in the steel sheets, and the mechanical properties and bending resistance of the steel sheets were excellent. Furthermore, among the steel sheets obtained in Example 1, those containing tempered martensite had a number density of precipitates in the range of 1 precipitate/ ^μm2 to 300 precipitates/ ^μm2 .

(Example 2)
Furthermore, in order to investigate the influence of the manufacturing conditions, hot-rolled sheets having a thickness of 2.8 mm were prepared at the hot-rolling finishing temperatures shown in Tables 5 to 7 for the steel types A to AZ for which excellent properties were recognized in Table 1. The hot-rolled sheets were wound, pickled, and cold-rolled to prepare cold-rolled sheets. The cold-rolled sheets were annealed and aged, and further optionally plated to obtain steel sheets for property evaluation. Here, the plated steel sheets were immersed in a hot-dip galvanizing bath and then held at the temperatures shown in Tables 5 to 7 to produce alloyed hot-dip galvanized steel sheets having an iron-zinc alloy plating layer on the surface of the steel sheets. In addition, for some steel sheets, a tempering treatment was performed in which the steel sheets were once cooled to 150 ° C. and reheated to a predetermined temperature and held for 2 seconds or more before the steel sheets were cooled to room temperature after being held at each residence temperature in the cold-rolled sheet annealing. The obtained results are shown in Tables 5 to 7. The evaluation method of the properties was the same as that in Example 1.

The results shown in Tables 5 to 7 reveal the following:

For A-2 and X-2, the aging time in aging treatment 1 was too short, resulting in a large difference in the amount of precipitates between aging treatment 1 and aging treatment 2. As a result, the difference in the number density of precipitates in the tempered martensite on the front and back of the steel plate became large, and the bending resistance of the steel plate decreased.

For C-2 and S-3, the annealing temperature during Q annealing was too high, which is believed to have caused the surface layer of the steel plate to be decarburized. As a result, the strength of the steel plate was reduced and sufficient bending resistance could not be secured.

For I-2 and W-2, the annealing temperature in IA annealing was too high (the annealing temperature was outside the range of the two-phase region of ferrite and austenite), so excessive austenite was produced, and the bending yield strength decreased due to the reduction in tempered martensite containing precipitates. In addition, it was not possible to secure 10.0% or more of retained austenite, and the elongation also decreased.

For L-2 and Z-3, the annealing temperature in IA annealing was too low (the annealing temperature was outside the range of the two-phase region of ferrite and austenite), so austenite was not sufficiently generated and the martensite obtained in Q annealing was tempered, making it impossible to ensure 10.0% or more of retained austenite. As a result, when bending deformation is applied to the steel plate, the effect of increasing strength due to the processing-induced transformation from retained austenite to martensite cannot be obtained, and sufficient bending resistance cannot be ensured.

For E-3 and AX-2, the annealing holding temperature in Q-annealing was low, and the base steel sheet could not be heated to the austenite single-phase region, so the martensite area ratio after Q-annealing was reduced. As a result, a sufficient amount of ε carbide could not be precipitated in the aging treatment, and sufficient bending resistance could not be ensured. In addition, the area ratio of retained austenite was reduced, so that it was not possible to obtain the effect of increasing strength due to the processing-induced transformation from retained austenite to martensite when bending deformation was applied to the steel sheet, and from this point of view, sufficient bending resistance could not be ensured.

For G-3 and U-2, it is believed that the bending radius in aging treatment 1 was too large, and therefore sufficient tensile strain was not generated to promote the formation of clusters or transition carbides. As a result, a sufficient amount of ε carbide could not be precipitated in aging treatment 1, and the difference in number density of precipitates in tempered martensite on the front and back of the steel plate became large, making it impossible to ensure sufficient bending resistance.

For M-3, it is believed that the bending radius in aging treatment 2 was too large, so sufficient tensile strain was not generated to promote the formation of clusters or transition carbides. As a result, a sufficient amount of ε carbide could not be precipitated in aging treatment 2, the difference in number density of precipitates in tempered martensite on the front and back of the steel sheet became large, and sufficient bending resistance could not be ensured.

For N-3 and AF-2, the aging time in aging treatment 2 was too short, resulting in a large difference in the amount of precipitates between aging treatment 1 and aging treatment 2. As a result, the difference in the number density of precipitates in the tempered martensite on the front and back of the steel plate became large, reducing the bending resistance of the steel plate.

For K-4 and AW-4, at least one of aging treatments 1 and 2 was omitted, so the number density of precipitates in the tempered martensite on at least one of the front and back sides of the steel plate was not controlled, and the difference in number density of precipitates in the tempered martensite on the front and back sides of the steel plate became large. As a result, the bending resistance of the steel plate was reduced.

In contrast, in the other examples, the steel sheets having the predetermined chemical composition were manufactured under the predetermined conditions, so that the predetermined metal structure was obtained in the steel sheets, and the steel sheets had excellent formability and bending resistance. In addition, among the steel sheets obtained in Example 2, those that were subjected to aging treatment and contained tempered martensite had a number density of precipitates in the range of ¹ precipitate/ ^μm2 to 300 precipitates/μm2.

From the above results, it can be said that steel plates satisfying the following requirements (I) to (IV) have excellent mechanical properties such as strength and elongation as well as excellent bending strength.

(I) In mass%, C: 0.10-0.30%, Si: 0.60-1.20%, Mn: 1.00-3.50%, P: 0.0200% or less, S: 0.0200% or less, Al: 0.001-1.000%, N: 0.0200% or less, Ti: 0-0.500%, Co: 0-0.500%, Ni: 0-0.500%, Mo: 0-0.500%, Cr: 0-2.000%, O: 0-0.0100%, B: 0-0.0100%, Nb: 0-0.500% , V: 0-0.500%, Cu: 0-0.500%, W: 0-0.1000%, Ta: 0-0.1000%, Sn: 0-0.0500%, Sb: 0-0.0500%, As: 0-0.0500%, Mg: 0-0.0500%, Ca: 0-0.0500%, Y: 0-0.0500%, Zr: 0-0.0500%, La: 0-0.0500%, and Ce: 0-0.0500%, with the balance being Fe and impurities.
(II) The steel structure includes, in terms of area ratio, a total of ferrite, pearlite and bainite: 0% or more and 30.0% or less, and retained austenite: 10.0% or more and 30.0% or less, with the remainder being martensite and tempered martensite.
(III) The difference between the number density of precipitates in tempered martensite on the first surface on the front side of the steel plate and the number density of precipitates in tempered martensite on the second surface on the back side of the steel plate is within 10.0%.
(IV) The yield strength is 600 MPa or more.

Furthermore, steel plate satisfying the above requirements (I) to (IV) can be manufactured by the following method.

hot rolling the steel slab having the chemical composition of (I) to obtain a hot rolled sheet;
coiling the hot-rolled sheet;
pickling the hot-rolled sheet;
cold rolling the hot rolled sheet to obtain a cold rolled sheet;
Q annealing is performed on the cold-rolled sheet;
The cold-rolled sheet subjected to the Q annealing is subjected to IA annealing; and
performing an aging treatment between the Q annealing and the IA annealing;
Including,
The Q annealing is a process of heating the cold-rolled sheet to a temperature in the austenite single phase region and 1000° C. or less, and then cooling the sheet to obtain a martensite structure having an area ratio of 90.0% or more.
The IA annealing is a process of holding the cold-rolled sheet in a two-phase region of ferrite and austenite to obtain retained austenite,
The aging treatment includes aging treatment 1 in which one of the front and back sides of the cold-rolled sheet is subjected to tensile deformation with a bending R of 2.0 m or less and held at 0 to 40 ° C. for 20 hr or more, and aging treatment 2 in which the other of the front and back sides of the cold-rolled sheet is subjected to tensile deformation with a bending R of 2.0 m or less and held at 0 to 40 ° C. for 20 hr or more.
Manufacturing method of steel plate.

Claims (6)

In mass percent,
C: 0.10-0.30%,
Si: 0.60-1.20%,
Mn: 1.00-3.50%,
P: 0.0200% or less,
S: 0.0200% or less,
Al: 0.001-1.000%,
N: 0.0200% or less,
Ti: 0 to 0.500%,
Co: 0 to 0.500%,
Ni: 0 to 0.500%,
Mo: 0-0.500%,
Cr: 0-2.000%,
O: 0 to 0.0100%,
B: 0 to 0.0100%,
Nb: 0 to 0.500%,
V: 0 to 0.500%,
Cu: 0-0.500%,
W: 0-0.1000%,
Ta: 0-0.1000%,
Sn: 0-0.0500%,
Sb: 0 to 0.0500%,
As: 0 to 0.0500%,
Mg: 0 to 0.0500%,
Ca: 0-0.0500%,
Y: 0 to 0.0500%,
Zr: 0 to 0.0500%,
La: 0 to 0.0500%, and
Ce: 0 to 0.0500%,
and the balance being Fe and impurities,
In terms of area ratio,
The sum of ferrite, pearlite and bainite: 0% or more and 30.0% or less, and
Retained austenite: 10.0% or more and 30.0% or less,
The balance has a steel structure consisting of fresh martensite and tempered martensite,
a difference between the number density of precipitates in the tempered martensite in the first surface on the front side of the steel sheet and the number density of precipitates in the tempered martensite in the second surface on the back side of the steel sheet is within 10.0%;
The yield strength is 600 MPa or more.
Steel plate. In mass percent,
Ti: 0.001 to 0.500%,
Co: 0.001 to 0.500%,
Ni: 0.001 to 0.500%,
Mo: 0.001-0.500%,
Cr:0.001~2.000%
O: 0.0001-0.0100%
B: 0.0001 to 0.0100%,
Nb: 0.001-0.500%,
V: 0.001-0.500%,
Cu: 0.001 to 0.500%,
W: 0.0001-0.1000%,
Ta: 0.0001 to 0.1000%,
Sn: 0.0001 to 0.0500%,
Sb: 0.0001 to 0.0500%,
As: 0.0001 to 0.0500%,
Mg: 0.0001-0.0500%,
Ca: 0.0001-0.0500%,
Y: 0.0001-0.0500%,
Zr: 0.0001 to 0.0500%,
La: 0.0001 to 0.0500%, and Ce: 0.0001 to 0.0500%,
The chemical composition contains one or more of the following:
The steel sheet according to claim 1. The steel structure contains the acicular retained austenite.
The steel sheet according to claim 1 or 2. The method for producing a steel sheet according to claim 1 or 2 ,
hot rolling a steel slab having the chemical composition according to claim 1 or 2 to obtain a hot rolled sheet;
coiling the hot-rolled sheet;
pickling the hot-rolled sheet;
cold rolling the hot rolled sheet to obtain a cold rolled sheet;
Q annealing is performed on the cold-rolled sheet;
The cold-rolled sheet subjected to the Q annealing is subjected to IA annealing; and
performing an aging treatment between the Q annealing and the IA annealing;
Including,
The Q annealing is a process of heating the cold-rolled sheet to a temperature in the austenite single phase region and 1000° C. or less, and then cooling the sheet to obtain a martensite structure having an area ratio of 90.0% or more.
The IA annealing is a process of holding the cold-rolled sheet in a two-phase region of ferrite and austenite to obtain retained austenite,
The aging treatment includes aging treatment 1 in which one of the front and back sides of the cold-rolled sheet is subjected to tensile deformation with a bending R of 2.0 m or less and held at 0 to 40 ° C. for 20 hr or more, and aging treatment 2 in which the other of the front and back sides of the cold-rolled sheet is subjected to tensile deformation with a bending R of 2.0 m or less and held at 0 to 40 ° C. for 20 hr or more.
Manufacturing method of steel plate. In the IA annealing, the cold-rolled sheet is held in a two-phase region of ferrite and austenite, and then cooled to room temperature. In the process, a coating layer made of zinc, aluminum, magnesium, or an alloy thereof is formed on the front and back surfaces of the cold-rolled sheet.
The method according to claim 4. In the IA annealing, needle-shaped retained austenite is obtained.
The method according to claim 4 or 5.