JP7323093B1 - High-strength steel plate and its manufacturing method - Google Patents

Abstract

An object of the present invention is to provide a high-strength steel sheet having a TS of 1180 MPa or more, a YR of 85% or more, excellent flatness in the sheet width direction and work embrittlement resistance, and a method for producing the same. It has a predetermined chemical composition, the amount of tempered martensite is 90% or more in area fraction, the amount of retained austenite is less than 3% in volume fraction, and the amount of ferrite and bainitic The total amount of ferrite is less than 10% in area fraction, the average grain size of prior austenite is 20 μm or less, and the average value of the occupancy of the packets having the maximum occupancy in the prior austenite grains is 70 in area fraction. % or less, high-strength steel plate.

Description

TECHNICAL FIELD The present invention relates to a high-strength steel sheet excellent in tensile strength, flatness in the sheet width direction, and resistance to work embrittlement, and a method for producing the same. The high-strength steel sheet of the present invention can be suitably used as structural members such as automobile parts.

With the aim of reducing _CO2 emissions by reducing the weight of vehicles and improving collision resistance by reducing the weight of vehicles, the strength of steel sheets for automobiles is being increased, and new regulations are being introduced one after another. . Therefore, for the purpose of increasing the strength of the vehicle body, the application of high-strength steel sheets with a tensile strength (TS) of 1180 MPa class or higher is increasing in the main structural parts forming automobiles.

High-strength steel sheets used in automobiles are required to have excellent work embrittlement resistance and an excellent yield ratio from the viewpoint of component performance. For example, in frame parts such as automobile bumpers, high-strength steel sheets with excellent work embrittlement resistance that does not cause embrittlement due to press forming and excellent impact absorption in the event of a collision, which is correlated with YR, are applied. It is preferable to

High-strength steel sheets used in automobiles are also required to have excellent flatness. Patent Literature 1 describes that warping of steel sheets adversely affects operational troubles in forming lines and dimensional accuracy of products. As a result of extensive studies, the inventors of the present invention have found that the dimensional accuracy of the product is affected not only by the warpage of the steel sheet, but also by the flatness in the width direction, which is evaluated using the steepness. For example, in order to achieve excellent dimensional accuracy, it is preferable that the steepness in the width direction is 0.02 or less.

In response to these demands, for example, Patent Document 2 provides a high-strength steel sheet having a tensile strength of 1100 MPa or more and excellent YR, surface properties and weldability, and a method for producing the same. However, the technique described in Patent Document 2 does not consider flatness in the plate width direction and work embrittlement resistance.

Patent Document 3 provides a hot-dip galvanized steel sheet having excellent press formability and low-temperature toughness and a tensile strength of 980 MPa or more, and a method for producing the same. However, although the technique described in Patent Document 3 can improve the embrittlement of the steel sheet due to the temperature drop, it does not consider the embrittlement of the steel sheet due to working. The flatness in the sheet width direction is not considered either.

Japanese Patent No. 4947176 Japanese Patent No. 6525114 Japanese Patent No. 6777272

Smart Process Society Journal 2013, Vol. 2, No. 3, p. 110-118

The present invention was developed in view of such circumstances, and provides a high-strength steel sheet having a TS of 1180 MPa or more, a YR of 85% or more, and excellent flatness in the width direction and work embrittlement resistance, and a method for producing the same. intended to

In order to achieve the above-described problems, the present inventors have made intensive studies and found the following.
(1) A TS of 1180 MPa or more can be achieved by setting the amount of tempered martensite to 90% or more.
(2) A YR of 85% or more can be achieved by setting the amount of retained austenite to less than 3% and the total amount of ferrite and bainitic ferrite to less than 10%.
(3) By setting the occupancy ratio of the maximum packet of tempered martensite to 70% or less in the prior austenite grains, excellent flatness in the sheet width direction can be achieved.
(4) Excellent work embrittlement resistance by setting the occupancy rate of the maximum packet of tempered martensite in the prior austenite grains to 70% or less and setting the average of the prior austenite grain size of the tempered martensite to 20 μm or less. characteristics can be realized.

The present invention has been made based on the above findings. That is, the gist and configuration of the present invention are as follows.
[1] % by mass, C: 0.030% or more and 0.500% or less, Si: 0.01% or more and 2.50% or less, Mn: 0.10% or more and 5.00% or less, P: 0.5% or less 100% or less, S: 0.0200% or less, Al: 1.000% or less, N: 0.0100% or less, and O: 0.0100% or less, the balance being Fe and unavoidable impurities In the component composition and the plate thickness 1/4 position, the amount of tempered martensite is 90% or more in area fraction, the amount of retained austenite is less than 3% in volume fraction, and the total amount of ferrite and bainitic ferrite is less than 10% in area fraction, the average prior austenite crystal grain size is 20 μm or less, and the average value of the occupancy of packets having the maximum occupancy in the prior austenite grains is 70% or less in area fraction. , high-strength steel plate.
[2] The component composition further includes, in mass%, Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.1% or less. 10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Co: 0.010% or less, Ni: 1.00% or less, Cu: 1.00% Below, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, The high-strength steel sheet according to [1], containing at least one element selected from Te: 0.100% or less, Hf: 0.10% or less, and Bi: 0.200% or less.
[3] The high-strength steel sheet according to [1] or [2], which has a plating layer on the surface of the steel sheet.
[4] The method for producing a high-strength steel sheet according to [1] or [2], wherein the steel having the chemical composition is subjected to hot rolling, pickling, and cold rolling to produce a cold-rolled sheet, The annealing temperature T1 is 750 ° C. or higher and 950 ° C. or lower, the holding time t1 at the annealing temperature T1 is 10 seconds or longer and 1000 seconds or shorter, and the average cooling rate from 750 ° C. to 600 ° C. is 20 ° C./s or higher, (Ms +50°C) to the quenching start temperature T2, the average cooling rate is 5°C/s or more and 30°C/s or less, and the quenching start temperature T2 is (Ms-80°C) or more and less than Ms, where Ms is in formula (1). The specified martensitic transformation start temperature (° C.), the average cooling rate from the rapid cooling start temperature T2 to 80° C. is cooled at 300° C./s or more, and the tempering temperature T3 is 100° C. or more and 400° C. or less. A method for producing a high-strength steel sheet, wherein the holding time t3 at T3 is 10 seconds or more and 10000 seconds or less.
Ms=519−474×[%C]−30.4×[%Mn]−12.1×[%Cr]−7.5×[%Mo]−17.7×[%Ni] (1)
Here, [%C], [%Mn], [%Cr], [%Mo], and [%Ni] represent the contents (% by mass) of C, Mn, Cr, Mo, and Ni, respectively, and do not include 0 if
[5] The method for producing a high-strength steel sheet according to [4], which includes plating.

According to the present invention, it is possible to obtain a high-strength steel sheet having a TS of 1180 MPa or more, a YR of 85% or more, and excellent flatness in the width direction and work embrittlement resistance. Further, by applying the high-strength steel sheet of the present invention to automobile structural members, for example, it is possible to improve fuel consumption by reducing the weight of the vehicle body. Therefore, the industrial utility value is extremely large.

FIG. 1 is a diagram showing the structure of a packet having the maximum occupation ratio in prior austenite grains according to the present invention and a calculation method thereof. FIG. 2 is a diagram showing the concept of the steepness θ of the steel sheet according to the present invention and its calculation method.

Embodiments of the present invention will be described below.

First, the appropriate range of the chemical composition of the high-strength steel sheet and the reason for its limitation will be described. In the following description, "%" representing the content of constituent elements of steel means "% by mass" unless otherwise specified.

[C: 0.030% or more and 0.500% or less]
C is one of the important basic components of steel, and is an important element that affects the fraction of tempered martensite and work embrittlement resistance in the present invention. If the C content is less than 0.030%, the fraction of tempered martensite decreases, making it difficult to achieve a TS of 1180 MPa or more. On the other hand, when the C content exceeds 0.500%, the tempered martensite becomes embrittled and the work embrittlement resistance is lowered. Therefore, the content of C is set to 0.030% or more and 0.500% or less. It is preferably 0.050% or more. It is preferably 0.400% or less. More preferably, it is 0.100% or more. More preferably, it is 0.350% or less.

[Si: 0.01% or more and 2.50% or less]
Si is one of the important basic components of steel, and in particular, in the present invention, it suppresses the formation of carbides during continuous annealing and promotes the formation of retained austenite. is an element. If the Si content is less than 0.01%, it becomes difficult to achieve a TS of 1180 MPa or more. On the other hand, when the Si content exceeds 2.50%, retained austenite increases excessively, making it difficult to achieve a YR of 85% or more. Therefore, the Si content should be 0.01% or more and 2.50% or less. Preferably, it is 0.05% or more. It is preferably 2.00% or less. More preferably, it is 0.10% or more. More preferably, it is 1.20% or less.

[Mn: 0.10% or more and 5.00% or less]
Mn is one of the important basic constituents of steel, and is an important element that particularly affects the fraction of tempered martensite and work embrittlement resistance in the present invention. If the Mn content is less than 0.10%, the fraction of tempered martensite decreases, making it difficult to achieve a TS of 1180 MPa or more. On the other hand, if the Mn content exceeds 5.00%, the tempered martensite becomes embrittled and the work embrittlement resistance is lowered. Therefore, the content of Mn is set to 0.10% or more and 5.00% or less. It is preferably 0.50% or more. It is preferably 4.50% or less. More preferably, it is 0.80% or more. More preferably, it is 4.00% or less.

[P: 0.100% or less]
Since P segregates at prior austenite grain boundaries and embrittles the grain boundaries, it lowers the ultimate deformability of the steel sheet, thereby lowering the work embrittlement resistance. Therefore, the P content should be 0.100% or less. Although the lower limit of the P content is not specified, it is preferably 0.001% or more because P is a solid-solution strengthening element and can increase the strength of the steel sheet. Therefore, the P content should be 0.100% or less. It is preferably 0.001% or more. It is preferably 0.070% or less.

[S: 0.0200% or less]
S exists as a sulfide and lowers the ultimate deformability of the steel sheet, thereby lowering the work embrittlement resistance. Therefore, the S content should be 0.0200% or less. Although the lower limit of the S content is not specified, it is preferably 0.0001% or more due to production technology restrictions. Therefore, the S content should be 0.0200% or less. It is preferably 0.0001% or more. It is preferably 0.0050% or less.

[Al: 1.000% or less]
Since Al raises the _A3 transformation point and contains a large amount of ferrite in the microstructure, the fraction of tempered martensite decreases, making it difficult to achieve a TS of 1180 MPa or more. Therefore, the Al content should be 1.000% or less. Although the lower limit of the Al content is not particularly specified, the Al content is preferably 0.001% or more because it suppresses the formation of carbides during continuous annealing and promotes the formation of retained austenite. Therefore, the Al content is set to 1.000% or less. It is preferably 0.001% or more. It is preferably 0.500% or less.

[N: 0.0100% or less]
N exists as a nitride and lowers the ultimate deformability of the steel sheet, thereby lowering the work embrittlement resistance. Therefore, the N content should be 0.0100% or less. Although the lower limit of the N content is not specified, it is preferable that the N content is 0.0001% or more due to production technology restrictions. Therefore, the content of N is set to 0.0100% or less. It is preferably 0.0001% or more. It is preferably 0.0050% or less.

[O: 0.0100% or less]
O exists as an oxide and lowers the ultimate deformability of the steel sheet, thereby lowering the work embrittlement resistance. Therefore, the O content should be 0.0100% or less. Although the lower limit of the O content is not particularly specified, it is preferable that the O content is 0.0001% or more due to production technology restrictions. Therefore, the O content is set to 0.0100% or less. It is preferably 0.0001% or more. It is preferably 0.0050% or less.

A high-strength steel sheet according to an embodiment of the present invention has a chemical composition containing the above components, with the balance being Fe and unavoidable impurities. Here, Zn, Pb, As, Ge, Sr and Cs are mentioned as unavoidable impurities. A total content of 0.100% or less of these impurities is allowed.

In addition to the above chemical composition, the high-strength steel sheet of the present invention further has, in mass%,
Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Ni: 1.00% or less, Co: 0.010% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.00% or less. 200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.10% and Bi: 0.200% or less may be contained singly or in combination.

If each of Ti, Nb and V is 0.200% or less, a large amount of coarse precipitates and inclusions are not generated, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the contents of Ti, Nb and V are each preferably 0.200% or less. Although the lower limits of the contents of Ti, Nb and V are not particularly specified, the strength of the steel sheet is increased by forming fine carbides, nitrides or carbonitrides during hot rolling or continuous annealing. Therefore, it is more preferable that the contents of Ti, Nb and V each be 0.001% or more. Therefore, when Ti, Nb and V are contained, their contents shall each be 0.200% or less. More preferably, it is 0.001% or more. More preferably, it is 0.100% or less.

If each of Ta and W is 0.10% or less, a large amount of coarse precipitates and inclusions are not generated, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, it is preferable that the contents of Ta and W are respectively 0.10% or less. Although the lower limits of the contents of Ta and W are not particularly specified, the strength of the steel sheet is increased by forming fine carbides, nitrides or carbonitrides during hot rolling or continuous annealing. More preferably, the contents of Ta and W are each 0.01% or more. Therefore, when Ta and W are contained, the content thereof should be 0.10% or less. More preferably, it is 0.01% or more. More preferably, it is 0.08% or less.

When B is 0.0100% or less, cracks are not generated inside the steel sheet during casting or hot rolling, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the B content is preferably 0.0100% or less. Although the lower limit of the B content is not particularly specified, it is an element that segregates at the austenite grain boundary during annealing and improves hardenability, so the content of B is preferably 0.0003% or more. preferable. Therefore, when B is contained, its content shall be 0.0100% or less. More preferably, it is 0.0003% or more. More preferably, it is 0.0080% or less.

When each of Cr, Mo and Ni is 1.00% or less, coarse precipitates and inclusions do not increase and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, it is preferable that the contents of Cr, Mo and Ni are each set to 1.00% or less. Although the lower limits of the contents of Cr, Mo, and Ni are not particularly specified, it is more preferable that the contents of Cr, Mo, and Ni are each 0.01% or more because they are elements that improve hardenability. . Therefore, when Cr, Mo and Ni are contained, their contents shall each be 1.00% or less. More preferably, it is 0.01% or more. More preferably, it should be 0.80% or less.

If Co is 0.010% or less, coarse precipitates and inclusions do not increase, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the Co content is preferably 0.010% or less. Although the lower limit of the Co content is not particularly specified, the Co content is more preferably 0.001% or more because it is an element that improves hardenability. Therefore, when Co is contained, its content should be 0.010% or less. More preferably, it is 0.001% or more. More preferably, it is 0.008% or less.

If Cu is 1.00% or less, coarse precipitates and inclusions do not increase, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the Cu content is preferably 1.00% or less. Although the lower limit of the Cu content is not particularly specified, the Cu content is preferably 0.01% or more because it is an element that improves hardenability. Therefore, when Cu is contained, its content shall be 1.00% or less. More preferably, it is 0.01% or more. More preferably, it should be 0.80% or less.

When Sn is 0.200% or less, cracks do not occur inside the steel sheet during casting or hot rolling, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the Sn content is preferably 0.200% or less. Although the lower limit of the Sn content is not particularly specified, the Sn content is set to 0.001% or more because Sn is an element that improves hardenability (generally, an element that improves corrosion resistance). is more preferable. Therefore, when Sn is contained, its content is made 0.200% or less. More preferably, it is 0.001% or more. More preferably, it is 0.100% or less.

If Sb is 0.200% or less, coarse precipitates and inclusions do not increase, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the Sb content is preferably 0.200% or less. Although the lower limit of the content of Sb is not particularly specified, the content of Sb is more preferably 0.001% or more because it is an element that controls the softened thickness of the surface layer and enables strength adjustment. Therefore, when Sb is contained, the content is made 0.200% or less. More preferably, it is 0.001% or more. More preferably, it is 0.100% or less.

When each of Ca, Mg and REM is 0.0100% or less, coarse precipitates and inclusions do not increase and the ultimate deformability of the steel sheet does not deteriorate, so the work embrittlement resistance does not deteriorate. Therefore, the contents of Ca, Mg and REM are preferably 0.0100% or less. Although the lower limits of the contents of Ca, Mg and REM are not particularly specified, since they are elements that spheroidize the shape of nitrides and sulfides and improve the ultimate deformability of the steel sheet, the contents of Ca, Mg and REM More preferably, the amount of each is 0.0005% or more. Therefore, when Ca, Mg and REM are contained, their contents shall each be 0.0100% or less. More preferably, it is 0.0005% or more. More preferably, it is 0.0050% or less.

If each of Zr and Te is 0.100% or less, coarse precipitates and inclusions do not increase and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the contents of Zr and Te are preferably 0.100% or less. Although the lower limits of the contents of Zr and Te are not particularly specified, the contents of Zr and Te are each 0 because they are elements that spheroidize the shape of nitrides and sulfides and improve the ultimate deformability of the steel sheet. It is more preferable to make it 0.001% or more. Therefore, when Zr and Te are contained, their contents shall each be 0.100% or less. More preferably, it is 0.001% or more. More preferably, it should be 0.080% or less.

If Hf is 0.10% or less, coarse precipitates and inclusions do not increase and the ultimate deformability of the steel sheet does not decrease, so the work embrittlement resistance does not decrease. Therefore, the Hf content is preferably 0.10% or less. Although the lower limit of the Hf content is not particularly specified, the Hf content is 0.01% or more because it is an element that spheroidizes the shape of nitrides and sulfides and improves the ultimate deformability of the steel sheet. is more preferable. Therefore, when Hf is contained, its content is made 0.10% or less. More preferably, it is 0.01% or more. More preferably, it is 0.08% or less.

If Bi is 0.200% or less, coarse precipitates and inclusions do not increase, and the ultimate deformability of the steel sheet is not lowered, so the work embrittlement resistance is not lowered. Therefore, the Bi content is preferably 0.200% or less. Although the lower limit of the Bi content is not particularly specified, the Bi content is more preferably 0.001% or more because it is an element that reduces segregation. Therefore, when Bi is contained, its content shall be 0.200% or less. More preferably, it is 0.001% or more. More preferably, it is 0.100% or less.

In addition, each content of Ti, Nb, V, Ta, W, B, Cr, Mo, Ni, Co, Cu, Sn, Sb, Ca, Mg, REM, Zr, Te, Hf and Bi is preferable. If it is less than the lower limit, it does not impair the effects of the present invention, so it is included as an unavoidable impurity.

Next, the steel structure of the high-strength steel sheet of the present invention will be explained.

[Tempered martensite: area fraction of 90% or more]
In the present invention, it is an extremely important invention constituent element. By using tempered martensite as the main phase, it is possible to achieve a TS of 1180 MPa or more. In order to obtain such an effect, the area fraction of tempered martensite must be 90% or more. Therefore, the area fraction of tempered martensite is set to 90% or more. Preferably it is 94% or more. More preferably, it is 96% or more.

Here, the method for measuring tempered martensite is as follows. After polishing the L cross section of the steel plate, 3 vol. % nital, and 1/4 part of the plate thickness (position corresponding to 1/4 of the plate thickness in the depth direction from the steel plate surface) is observed in 10 fields of view at a magnification of 2000 using SEM. In the above structure image, tempered martensite is a structure in which the inside of the structure has fine irregularities and which has carbide inside. The tempered martensite can be determined from the average of these values.

[Amount of retained austenite is less than 3%]
In the present invention, it is an extremely important invention constituent element. If the volume fraction of retained austenite is 3% or more, it becomes difficult to achieve a YR of 85% or more. The cause of the decrease in YR is that an increase in retained austenite causes a decrease in YS due to deformation-induced transformation of retained austenite. Therefore, retained austenite should be less than 3%. It is preferably 1% or less. In addition, the lower limit of retained austenite is not particularly limited. It may be 0%.

Here, the method for measuring retained austenite is as follows. Retained austenite is obtained by polishing the steel plate from 1/4 part of the plate thickness to a surface of 0.1 mm, and then chemically polishing the surface to 0.1 mm. }, {220}, {311} planes and {200}, {211}, {220} planes of bcc iron. Measure the integrated intensity ratio of each diffraction peak, and average the nine integrated intensity ratios obtained. and ask for it.

[Total of ferrite and bainitic ferrite: less than 10% in area fraction]
In the present invention, it is an extremely important invention constituent element. If the total content of ferrite and bainitic ferrite is 10% or more, it becomes difficult to achieve a TS of 1180 MPa or more and a YR of 85% or more. The cause of the decrease in YR is that ferrite and bainitic ferrite are soft structures and yield early. Therefore, the total content of ferrite and bainitic ferrite should be less than 10%. It is preferably 8% or less. More preferably, it is 5% or less. The lower limit of the sum of ferrite and bainitic ferrite is not particularly limited. It may be 0%.

Here, the method for measuring the sum of ferrite and bainitic ferrite is as follows. After polishing the L cross section of the steel plate, 3 vol. % nital, and 1/4 part of the plate thickness (position corresponding to 1/4 of the plate thickness in the depth direction from the steel plate surface) is observed in 10 fields of view at a magnification of 2000 using SEM. In the above structure images, the ferrite and the bainitic ferrite are concave structures with flat interiors and no carbides inside. The sum of ferrite and bainitic ferrite can be obtained from the average of those values.

Perlite, fresh martensite, acicular ferrite, and the like are conceivable as structures other than the whole structure described above. These structures do not affect the characteristics as long as they are in the range of 5% or less, so they may be included.

[Prior austenite average crystal grain size is 20 μm or less]
In the present invention, it is an extremely important invention constituent element. By reducing the average grain size of prior austenite, propagation of cracks can be suppressed, so that the work embrittlement resistance of the steel sheet is improved. In order to obtain such an effect, the prior austenite average crystal grain size must be 20 μm or less. Although the lower limit of the prior-austenite average crystal grain size is not particularly specified, if the prior-austenite average crystal grain size is less than 2 μm, retained austenite may increase, so it is preferably 2 μm or more. Therefore, the average grain size of prior austenite is set to 20 μm or less. The thickness is preferably 2 μm or more. The thickness is preferably 15 μm or less. More preferably, the thickness is 3 μm or more. More preferably, the thickness is 10 μm or less.

Here, the method for measuring the prior austenite average crystal grain size is as follows. After polishing the L cross section of the steel sheet, it is etched with a mixed solution of picric acid and ferric chloride, etc., and after exposing the prior austenite grain boundaries, 1/4 part of the sheet thickness (from the steel sheet surface in the depth direction 3 to 10 fields of view are photographed with an optical microscope at a magnification of 400 times. A total of 20 straight lines of 10 vertical×10 horizontal are drawn at equal intervals on the obtained image data and obtained by the cutting method.

[Average value of occupancy of packets having maximum occupancy in prior austenite grains is 70% or less in terms of area fraction]
In the present invention, it is an extremely important invention constituent element. The occupancy of the packet having the maximum occupancy within the prior austenite grains affects the flatness in the sheet width direction and the resistance to work embrittlement. The packet having the maximum occupancy rate in the prior austenite grains is, as shown in FIG. indicates the packet with the largest occupancy of . The occupancy of one packet within a prior austenite grain is determined by dividing the area of the specified packet by the total area within the prior austenite grain. As a result of extensive studies, the present inventors found that by reducing the occupancy ratio of packets having the maximum occupancy ratio in the prior austenite grains, the strain between packets was alleviated and the flatness in the sheet width direction was improved. I found out. In addition, it was found that reducing the occupancy of the packet, which has the maximum occupancy in the prior austenite grains, refines the structure and suppresses the propagation of cracks, thereby improving the work embrittlement resistance of the steel sheet. . Therefore, the average value of the occupancy of the packets having the maximum occupancy within the prior austenite grains is set to 70% or less. It is preferably 60% or less. In addition, the lower limit of the average value of the occupancy ratio of the packet having the maximum occupancy ratio in the prior austenite grains is not particularly limited. The maximum number of packet types is four, and when the four packets are evenly distributed, the packet having the maximum occupation rate in the old austenite grains has a occupancy rate of 25%. Therefore, the lower limit of the average value of the occupancy ratio of the packets having the maximum occupancy ratio in the prior austenite grains is 25% or more, but it is not necessary to be limited to this.

Here, the method for measuring the average value of the occupancy ratio of the packets having the maximum occupancy ratio within the prior austenite grains is as follows. First, a test piece for structure observation is taken from a cold-rolled steel sheet. Next, the sampled test piece is polished by colloidal silica vibration polishing so that the cross section in the rolling direction (L cross section) serves as the observation surface. The observation surface shall be a mirror surface. Next, electron beam backscatter diffraction (EBSD) measurement is performed on 1/4 part of the plate thickness (position corresponding to 1/4 of the plate thickness in the depth direction from the steel plate surface) to obtain local crystal orientation data. At this time, the SEM magnification is 1000 times, the step size is 0.2 μm, the measurement area is 80 μm square, and the WD is 15 mm. The obtained local orientation data is analyzed using OIMAnalysis 7 (OIM), and a diagram (CP map) colored for each Close-packed Plane group (CP group) is created using the method described in Non-Patent Document 1. . In the present invention, a packet is defined as an area belonging to the same CP group. From the obtained CP map, the area of the packet having the largest occupancy is obtained and divided by the total area within the prior austenite grains to obtain the occupancy of the packet having the maximum occupancy within the prior austenite grains. This analysis is performed for 10 or more adjacent prior austenite grains, and the average value is taken as the average value of the occupancy of the packet having the maximum occupancy within the prior austenite grains.

Next, the manufacturing method of the present invention will be described.

In the present invention, the method of melting the steel material (steel slab) is not particularly limited, and any known melting method such as a converter or an electric furnace is suitable. Steel slabs (slabs) are preferably produced by continuous casting to prevent macro-segregation.

In the present invention, the slab heating temperature, slab soaking holding time and coiling temperature in hot rolling are not particularly limited. Methods for hot rolling steel slabs include a method of rolling after heating the slab, a method of directly rolling the slab after continuous casting without heating, and a method of subjecting the slab after continuous casting to heat treatment for a short period of time before rolling. etc. The slab heating temperature, slab soaking holding time, finish rolling temperature and coiling temperature in hot rolling are not particularly limited, but the lower limit of the slab heating temperature is preferably 1100° C. or higher. The upper limit of the slab heating temperature is preferably 1300°C or less. The lower limit of the slab soaking holding time is preferably 30 minutes or longer. The upper limit of the slab soaking holding time is preferably 250 minutes or less. The lower limit of the finish rolling temperature is preferably at least the Ar ₃ transformation point. Moreover, the lower limit of the winding temperature is preferably 350° C. or higher. Moreover, the upper limit of the winding temperature is preferably 650° C. or less.

The hot-rolled steel sheet thus produced is pickled. Since pickling can remove oxides from the surface of the steel sheet, it is important for ensuring good chemical conversion treatability and plating quality in the final high-strength steel sheet. Also, the pickling may be performed once, or may be divided into a plurality of times. Further, the hot-rolled pickling-treated sheet may be cold-rolled, or the cold-rolled sheet may be heat-treated and then cold-rolled.

The rolling reduction in cold rolling and the plate thickness after rolling are not particularly limited, but the lower limit of the rolling reduction is preferably 30% or more. Moreover, the upper limit of the rolling reduction is preferably 80% or less. The number of rolling passes and the rolling reduction of each pass are not particularly limited, and the effects of the present invention can be obtained.

The cold-rolled steel sheet obtained as described above is annealed. Annealing conditions are as follows.

[Annealing temperature T1: 750°C or higher and 950°C or lower]
When the annealing temperature T1 is less than 750°C, the total area fraction of ferrite and bainitic ferrite is 10% or more, making it difficult to achieve a TS of 1180 MPa or more and a YR of 85% or more. become difficult to do. On the other hand, if the annealing temperature T1 exceeds 950° C., the prior austenite grain size increases excessively, exceeding 20 μm, and the work embrittlement resistance deteriorates. Therefore, the annealing temperature T1 is set at 750°C or higher and 950°C or lower. It is preferably 800° C. or higher. It is preferably 900° C. or less.

[Holding time t1 at annealing temperature T1: 10 seconds or more and 1000 seconds or less]
When the holding time t1 at the annealing temperature T1 is less than 10 seconds, the austenitization is insufficient, and the total area fraction of ferrite and bainitic ferrite is 10% or more, making it difficult to achieve a TS of 1180 MPa or more. and it becomes difficult to realize YR of 85% or more. On the other hand, when the holding time at the annealing temperature T1 exceeds 1000 seconds, the grain size of the prior austenite is excessively increased, and the work embrittlement resistance is deteriorated. Therefore, the holding time t1 at the annealing temperature T1 is set to 10 seconds or more and 1000 seconds or less. It is preferably 50 seconds or more. It is preferably 500 seconds or less.

[Average cooling rate from 750°C to 600°C is 20°C/s or more]
If the average cooling rate from 750 ° C. to 600 ° C. is less than 20 ° C./s, the total area fraction of ferrite and bainitic ferrite is 10% or more, making it difficult to achieve a TS of 1180 MPa or more, and , it becomes difficult to achieve a YR of 85% or more. Therefore, the average cooling rate from 750°C to 600°C should be 20°C/s or more. It is preferably 30° C./s or more. The upper limit is not particularly limited, but is preferably 2000° C./s or less.

[Average cooling rate from (Ms + 50 ° C.) to rapid cooling start temperature T2 is 5 ° C./s or more and 30 ° C./s or less]
In the present invention, it is an extremely important invention constituent element. The average cooling rate from (Ms+50° C.) to the quenching start temperature T2 affects the average value of the total area fraction of ferrite and bainitic ferrite and the average occupancy of packets having the maximum occupancy in prior austenite grains. When the average cooling rate from (Ms + 50°C) to the rapid cooling start temperature T2 is less than 5°C/s, the total area fraction of ferrite and bainitic ferrite is 10% or more, making it difficult to achieve a TS of 1180 MPa or more. and it becomes difficult to realize YR of 85% or more. On the other hand, when the average cooling rate from (Ms + 50 ° C.) to the quenching start temperature T2 exceeds 30 ° C./s, the average value of the occupancy of the packet having the maximum occupancy in the prior austenite grains exceeds 70%, and the width direction flatness deteriorates, and work embrittlement resistance deteriorates. Therefore, the average cooling rate from (Ms+50° C.) to the rapid cooling start temperature T2 should be 5° C./s or more and 30° C./s or less. It is preferably 10° C./s or more. It is preferably 20° C./s or less.

[Rapid cooling start temperature T2 is (Ms - 80 ° C.) or more and less than Ms]
In the present invention, it is an extremely important invention constituent element. The quenching start temperature T2 is set to (Ms−80° C.) or more and less than Ms, and quenching is performed in a state where the martensite transformation rate before the start of quenching is 1% or more and 80% or less. As a result, it is possible to obtain a structure in which the average value of the occupancy of packets having the maximum occupancy in the prior austenite grains is 70% or less and the volume fraction of retained austenite is less than 3%. When the quenching start temperature T2 is less than (Ms-80 ° C.), the martensitic transformation rate before the start of quenching exceeds 80%, so the retained austenite is 3% or more in volume fraction, realizing a YR of 85% or more. become difficult to do. On the other hand, when the quenching start temperature T2 exceeds Ms, the martensitic transformation rate before the start of quenching is less than 1%, and the average value of the packet occupancy rate having the maximum occupancy rate in the prior austenite grains exceeds 70%. The flatness in the width direction deteriorates, and the work embrittlement resistance deteriorates. Therefore, the rapid cooling start temperature T2 is set at (Ms−80° C.) or more and less than Ms. It is preferably (Ms-50°C) or more. It is preferably (Ms-5°C) or less. The martensitic transformation start temperature Ms (°C) is defined by the following formula (1).
Ms=519−474×[%C]−30.4×[%Mn]−12.1×[%Cr]−7.5×[%Mo]−17.7×[%Ni] (1)
Here, [%C], [%Mn], [%Cr], [%Mo], and [%Ni] represent the contents (% by mass) of C, Mn, Cr, Mo, and Ni, respectively, and do not include 0 if

[Average cooling rate from quenching start temperature T2 to 80°C is 300°C/s or more]
If the average cooling rate from the rapid cooling start temperature T2 to 80°C is less than 300°C/s, the volume fraction of retained austenite is 3% or more, making it difficult to achieve a YR of 85% or more. Therefore, the average cooling rate from the rapid cooling start temperature T2 to 80°C is set to 300°C/s or more. It is preferably 800° C./s or higher. The upper limit is not particularly limited, but is preferably 2000° C./s or less.

[Tempering temperature T3 is 100°C or higher and 400°C or lower]
In the present invention, tempered martensite refers to a structure obtained by heat-treating martensite at a temperature of 80° C. or lower at a tempering temperature of 100° C. or higher for a holding time of 10 seconds or longer. Therefore, when the tempering temperature T3 is less than 100° C., the martensite is not sufficiently tempered, and the as-quenched martensite structure becomes the main structure, and the as-quenched martensite deteriorates the work embrittlement resistance. On the other hand, when the tempering temperature T3 exceeds 400°C, the tempered martensite decomposes and becomes ferrite, so the area fraction of the tempered martensite becomes less than 90%, making it difficult to achieve a TS of 1180 MPa or more. Become. Therefore, the tempering temperature T3 is set to 100°C or higher and 400°C or lower. The temperature is preferably 150° C. or higher. The temperature is preferably 350° C. or lower.

[Holding time t3 at tempering temperature T3 is 10 seconds or more and 10000 seconds or less]
In the present invention, tempered martensite refers to a structure obtained by heat-treating martensite at a temperature of 80° C. or lower at a tempering temperature of 100° C. or higher for a holding time of 10 seconds or longer. Therefore, if the holding time t3 at the tempering temperature T3 is less than 10 seconds, the martensite is not sufficiently tempered, and the as-quenched martensite structure becomes the main structure, and the as-quenched martensite deteriorates the work embrittlement resistance. On the other hand, when the tempering temperature T3 exceeds 10000 seconds, the tempered martensite decomposes and becomes ferrite, so the area fraction of the tempered martensite becomes less than 90%, making it difficult to achieve a TS of 1180 MPa or more. Become. Therefore, the holding time t3 at the tempering temperature T3 is set to 10 seconds or more and 10000 seconds or less. It is preferably 50 seconds or more. It is preferably 5000 seconds or less.

Cooling after tempering does not have to be specified, and it may be cooled to a desired temperature by any method. It should be noted that the desired temperature is desirably about room temperature.

Further, the above high-strength steel sheet may be processed under the condition that the equivalent plastic strain amount is 0.10% or more and 5.00% or less. Further, after processing, reheating may be performed under the condition of 100° C. or more and 400° C. or less.

When high-strength steel sheets are traded, they are usually traded after being cooled to room temperature.

The high-strength steel sheet may be plated during or after annealing.

Examples of the plating treatment during annealing include hot dip galvanizing treatment during or after cooling at an average cooling rate of 20° C./s or more from 750° C. to 600° C. and alloying treatment after hot dip galvanizing. As the plating treatment after annealing, for example, Zn—Ni electro-alloy plating treatment after tempering or pure Zn electroplating treatment can be exemplified. A plating layer may be formed by electroplating, or hot-dip zinc-aluminum-magnesium alloy plating may be applied. In addition, in the above-mentioned plating treatment, the case of zinc plating was mainly explained, but the type of plating metal such as Zn plating and Al plating is not particularly limited. Other manufacturing method conditions are not particularly limited, but from the viewpoint of productivity, a series of treatments such as the above annealing, hot-dip galvanizing, galvanizing treatment, etc. Line). After hot-dip galvanization, wiping is possible in order to adjust the basis weight of the plating. In addition, the conditions of plating etc. other than the above-mentioned conditions can be based on the usual method of hot-dip galvanization.

After plating after annealing, processing may be performed again under the condition that the equivalent plastic strain amount is 0.10% or more and 5.00 or less. Further, after processing, reheating may be performed under the condition of 100° C. or more and 400° C. or less.

A steel having the chemical composition shown in Tables 1 and 2, with the balance being Fe and unavoidable impurities, was melted in a converter and made into a slab by a continuous casting method. Next, the obtained slab was heated, subjected to pickling treatment after hot rolling, cold rolling, annealing treatment and tempering treatment shown in Tables 3 to 6, and the plate thickness was 0.6. A high-strength cold-rolled steel sheet of ~2.2 mm was obtained. Some of the steel sheets are manufactured by plating during or after annealing.

Using the high-strength cold-rolled steel sheets obtained as described above as test steels, tensile properties, flatness in the sheet width direction, and resistance to work embrittlement were evaluated according to the following test methods.

(Organization observation)
According to the method described above, the total amount of tempered martensite, retained austenite, ferrite, bainitic ferrite, and the average grain size of prior austenite were determined.

(Occupancy of packet with maximum occupancy in prior austenite grains)
According to the method described above, the average value of the occupancy of the packet having the maximum occupancy within the prior austenite grains was obtained.

(Tensile test)
For the tensile test, a JIS No. 5 test piece (gauge length: 50 mm, width of parallel part: 25 mm) was sampled so that the longitudinal direction of the test piece was perpendicular to the rolling direction, and the test was performed according to JIS Z 2241. A tensile test was performed at a crosshead speed of 1.67×10 ⁻¹ mm/sec to measure YS and TS. In the present invention, a TS of 1180 MPa or more was judged to be acceptable. A yield ratio (YR) of 85% or more was judged to be acceptable. Note that YR is obtained by the following equation (2).
YR=100×YS/TS (2)
(flatness in width direction)
For the various cold-rolled steel sheets obtained as described above, the flatness in the sheet width direction was measured by the method shown in FIG. Specifically, a plate (coil width x 500 mmL x plate thickness) with a length of 500 mm in the rolling direction is cut out from the coil, placed on the surface plate so that the warp of the end face faces upward, and the stylus is placed on the object to be measured. The height of the steel plate was continuously measured over the entire width direction using a moving contact-type displacement gauge. Based on the results, the steepness θ, which is an index indicating the flatness of the shape of the steel sheet, was measured according to the method shown in FIG. A steepness exceeding 0.02 was evaluated as "x", a steepness exceeding 0.01 and 0.02 or less was evaluated as "○", and a steepness of 0.01 or less was evaluated as "◎". A flatness of 0.02 or less was judged to be "excellent in flatness in the plate width direction".

(Processing embrittlement resistance)
The resistance to working brittleness was evaluated by the Charpy test. For the Charpy test piece, a plurality of steel plates were superimposed and fastened with bolts, and after confirming that there was no gap between the steel plates, a test piece with a V-notch having a depth of 2 mm was produced. The number of steel sheets to be laminated was set so that the thickness of the test piece after lamination was closest to 10 mm. For example, when the plate thickness is 1.2 mm, 8 sheets are laminated and the thickness of the test piece becomes 9.6 mm. A laminated Charpy test piece was taken with the sheet width direction as the longitudinal side. As an index showing the resistance to work embrittlement, the ratio vE _0% /vE _10% of impact absorption energy at room temperature between the as-manufactured (unworked) steel sheet and the steel sheet subjected to 10% rolling was measured. "X" for vE _0% / vE _10% less than 0.6, "○" for vE _0% / vE _10% of 0.6 or more and less than 0.7, vE _0% / vE _10% Those having a vE _0% /vE 10% ratio of 0.7 or more were evaluated as "excellent", and those having a vE 0% /vE _10% ratio of 0.6 or more were judged to be "excellent in working embrittlement resistance". Conditions other than the above conformed to JIS Z 2242:2018.

The results are shown in Tables 7-10. As shown in these tables, the examples of the present invention have a TS of 1180 MPa or more, a YR of 85% or more, and are excellent in flatness in the width direction and resistance to work embrittlement. On the other hand, the comparative examples are inferior in one or more of TS, YR, flatness in the sheet width direction, and work embrittlement resistance.

Claims (5)

in % by mass,
C: 0.030% or more and 0.500% or less,
Si: 0.01% or more and 2.50% or less,
Mn: 0.10% or more and 5.00% or less,
P: 0.100% or less,
S: 0.0200% or less,
Al: 1.000% or less,
N: 0.0100% or less, and
A component composition containing O: 0.0100% or less, the balance being Fe and inevitable impurities,
At the plate thickness 1/4 position,
The amount of tempered martensite is 90% or more in terms of area fraction,
The amount of retained austenite is less than 3% in volume fraction,
The total area fraction of ferrite and bainitic ferrite is less than 10%,
The prior austenite average crystal grain size is 20 μm or less,
A high-strength steel sheet, wherein the average value of the occupancy of packets having the maximum occupancy in prior austenite grains is 70% or less in terms of area fraction. The component composition further, in mass %,
Ti: 0.200% or less, Nb: 0.200% or less,
V: 0.200% or less, Ta: 0.10% or less,
W: 0.10% or less, B: 0.0100% or less,
Cr: 1.00% or less, Mo: 1.00% or less,
Co: 0.010% or less, Ni: 1.00% or less,
Cu: 1.00% or less, Sn: 0.200% or less,
Sb: 0.200% or less, Ca: 0.0100% or less,
Mg: 0.0100% or less, REM: 0.0100% or less,
Zr: 0.100% or less, Te: 0.100% or less,
Hf: 0.10% or less, Bi: 0.200% or less,
The high-strength steel sheet according to claim 1, containing at least one element selected from The high-strength steel sheet according to claim 1 or 2, which has a plating layer on the surface of the steel sheet. A method for manufacturing a high-strength steel sheet according to claim 1 or 2,
A cold-rolled sheet produced by hot-rolling, pickling and cold-rolling the steel having the above composition,
Annealing temperature T1 is 750° C. or higher and 950° C. or lower,
Heating under the condition that the holding time t1 at the annealing temperature T1 is 10 seconds or more and 1000 seconds or less,
An average cooling rate of 20°C/s or more from 750°C to 600°C,
The average cooling rate from (Ms + 50 ° C.) to the rapid cooling start temperature T2 is 5 ° C./s or more and 30 ° C./s or less,
The quenching start temperature T2 is (Ms−80° C.) or more and less than Ms, where Ms is the martensite transformation start temperature (° C.) defined by formula (1),
Cooling at an average cooling rate of 300 ° C./s or more from the quenching start temperature T2 to 80 ° C.,
Tempering temperature T3 is 100°C or higher and 400°C or lower,
A method for producing a high-strength steel sheet, comprising heating at a tempering temperature T3 for a holding time t3 of 10 seconds or more and 10000 seconds or less.
Ms=519−474×[%C]−30.4×[%Mn]−12.1×[%Cr]−7.5×[%Mo]−17.7×[%Ni] (1)
Here, [%C], [%Mn], [%Cr], [%Mo], and [%Ni] represent the contents (% by mass) of C, Mn, Cr, Mo, and Ni, respectively, and do not include 0 if The method for producing a high-strength steel sheet according to claim 4, wherein plating is applied.

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