JP7092264B2 - High-strength steel plate - Google Patents

Description

The present invention relates to a high-strength steel sheet.
The present application claims priority based on Japanese Patent Application No. 2019-112081 filed in Japan on June 17, 2019, the contents of which are incorporated herein by reference.

In recent years, the ratio of high-strength steel sheets used in automobile parts has been increasing for the purpose of improving the fuel efficiency and collision safety of automobiles. Since steel sheets tend to deteriorate in workability as the strength increases, conventionally, DP (Dual Phase) steel and TRIP (Transformation) have been used.
Induced Plasticity) Steel sheets such as steel that can achieve both high strength and workability have been developed. Further, from the viewpoint of further enhancing the collision safety of automobiles, it is also required for automobile parts to increase the absorbed energy at the time of collision. To meet this demand, it is effective to increase the yield strength of the steel sheets that make up the automobile parts. However, the higher the yield strength of the steel sheet, the more likely it is that the ductility will decrease. When a steel sheet having a high yield strength and a low ductility is used for an automobile part, the steel sheet may be broken due to insufficient elongation at the time of a collision of an automobile. When the steel sheet is broken in this way, the steel sheet cannot sufficiently absorb the impact and may not be able to sufficiently exhibit its performance due to its high yield strength.

Automotive parts may be manufactured by press forming steel plates. In the painting process (painting baking process) after assembling the automobile parts obtained by press-molding the steel sheet into the car body skeleton, seizure hardening occurs on the steel sheet. Seizure hardening is a phenomenon in which the yield strength of a steel sheet increases due to strain aging, and is also called strain aging hardening. The seizure hardening is caused by an intrusive element such as a solid solution C or a solid solution N in the steel. Since DP steel and TRIP steel contain a large amount of C, the amount of shrinkage hardening (the amount of increase in the yield strength of the steel sheet due to strain aging) is large. Therefore, the ductility of conventional automobile parts made of DP steel or TRIP steel may decrease even if the yield strength increases due to seizure hardening. To solve such problems, in order to secure ductility after baking and hardening, a technique for reducing hard phases such as martensite and bainite in the steel structure while ensuring the strength of the steel sheet is being studied.

For example, Patent Document 1 discloses a steel sheet in which the hard phase such as bainite and martensite is suppressed to a small amount. The steel sheet described in Patent Document 1 has ferrite as the main phase, contains 2 to 12% of pearlite and 3% or less of martensite in volume fraction, and the balance is composed of a low temperature generated phase.

International Publication No. 2017/169941

However, the steel sheet described in Patent Document 1 has a tensile strength of 440 MPa class, and its strength is insufficient to be applied to a shock absorbing member used in recent automobiles and the like. Further, Patent Document 1 does not mention the ductility of a steel sheet when coating and baking are performed after press forming.

As described above, the conventional high-strength steel sheet may tend to have a decrease in ductility as the yield strength increases due to seizure hardening. In view of the above circumstances, it is an object of the present invention to provide a high-strength steel sheet in which a decrease in ductility due to baking hardening is suppressed.

As a result of repeated studies on the ductility of the steel sheet after the coating baking treatment, the present inventors obtained the following findings (a) and (b) from the experimental results.
(A) A predetermined amount of Ni and Al are simultaneously contained to form a columnar precipitation phase of Ni—Al (an intermetallic compound of Ni—Al), and the main phase is a structure composed of ferrite, cementite and / or pearlite. The steel plate containing a small amount of the hard phase of bainite and / or martensite has a tensile strength of 590 MPa or more, and the amount of change in elongation in the tensile test is BH-treated (after 2% prestrain is applied, 20 at 170 ° C.). It will be 4% or less before and after (heat treatment for minutes). (B) In order to obtain a steel sheet having a steel structure as described above, it is important to control the cooling rate after hot rolling, and the subsequent quenching conditions and winding temperature.

The reason why the ductility of the above steel sheet after the coating baking treatment is good is not clear, but the present inventors presume as follows. Strain aging hardening (baking hardening) is a phenomenon in which carbon adheres to dislocations introduced by processing to stabilize dislocations and increase resistance to dislocation motion, thereby increasing yield strength. The phase transformation from austenite to martensite or bainite involves the formation of dislocations. Therefore, when compared with steel sheets of similar strength, the dislocation density of the steel sheet according to the present invention containing a small amount of bainite and / or martensite is lower than that of the conventional steel sheet containing a large amount of martensite and / or bainite. .. Similarly, even after the prestrain is applied, the dislocation density of the steel sheet according to the present invention is lower than the dislocation density of the conventional steel sheet. For these reasons, it can be inferred that the steel sheet according to the present invention having a low structure fraction of bainite and / or martensite has a small amount of dislocations stabilized by strain age hardening and suppresses a decrease in ductility.

The present invention has been made based on the above findings, and the gist thereof is as follows. [1] The high-strength steel sheet according to one aspect of the present invention has a component composition of mass%.
C: 0.0150% or more and 0.3000% or less,
Si: More than 0% and less than 3.000%,
Mn: 0.050% or more and 3.600% or less,
P: More than 0% and 0.030% or less,
S: More than 0% and 0.0200% or less,
Al: 0.500% or more and 5.000% or less,
N: More than 0% and less than 0.0100%,
Ni: 1.000% or more and 12.300% or less,
Cu: 0% or more and 4.800% or less,
Mo: 0% or more and 2.500% or less,
Ca: 0% or more and 0.0200% or less,
Mg: 0% or more and 0.0200% or less, and REM: 0% or more and 0.0200% or less, and the balance consists of Fe and impurities.
Satisfying the relationships of the following equations (1), (2) and (3) at the same time,
Steel structure from ferrite, cementite and pearlite totaling 85.0% or more and 99.0% or less, and martensite, bainite and retained austenite totaling 1.0% or more and 15.0% or less in area ratio. Configured,
The columnar precipitation phase of Ni—Al is present on the grain boundaries of ferrite, and the coverage of the ferrite grain boundaries by the columnar precipitation phase is 5.0% or more of the total grain boundary length.
The tensile strength is 590 MPa or more.
(Ni% + 12.00 x C% + 2.00 x Mn% + 1.20 x Cu%) -2.00 x (Al% + 0.50 x Si% + 0.25 x Mo%) ≧ 0.0 (1)
3.725 x C% + 0.160 x Si% + 0.630 x Mn% -0.110 x Al% + 0.210 x Ni% + 0.450 x Cu% + 0.620 x Mo% -1.818 ≦ 1. 000
(2)
Ni% -0.5 x Cu% ≧ 0.0 (3)
Here, C%, Si%, Mn%, Al%, Ni%, Cu% and Mo% are the contents of C, Si, Mn, Al, Ni, Cu and Mo in mass%, respectively, and are contained. If not, substitute 0.
[2] The high-strength steel sheet according to the above [1] has a component composition of mass%.
Cu: 0.050% or more and 4.800% or less,
Mo: 0.030% or more and 2.500% or less,
Ca: 0.0001% or more and 0.0200% or less,
It may contain at least one selected from Mg: 0.0001% or more and 0.0200% or less, and REM: 0.0001% or more and 0.0200% or less.
[3] The high-strength steel sheet according to the above [1] or [2] may have a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface.

According to the above aspect according to the present invention, it is possible to provide a high-strength steel sheet in which a decrease in ductility due to baking hardening is suppressed. The high-strength steel sheet according to the above aspect is suitable for a shock absorbing member that requires collision resistance.

Hereinafter, a high-strength steel sheet according to an embodiment of the present invention (hereinafter, may be referred to as a high-strength steel sheet or simply a steel sheet according to the present embodiment) will be described. However, the present invention is not limited to the configuration disclosed in the present embodiment, and various modifications can be made without departing from the spirit of the present invention.
The high-strength steel sheet according to this embodiment has a component composition (chemical composition) and a metal structure described below. In the present embodiment, the "high-strength steel sheet" means a steel sheet having a tensile strength (TS) of 590 MPa or more.

(Ingredient composition)
First, the composition of the high-strength steel sheet according to the present embodiment will be described. All "%" for the content of each element mean "mass%".

"C: 0.0150% or more and 0.3000% or less"
C is an essential element for forming carbides in steel to obtain bainite, martensite and retained austenite. If the C content is less than 0.0150%, the total area ratio of martensite, bainite and retained austenite cannot be 1.0% or more. Therefore, the C content is set to 0.0150% or more.
On the other hand, when the C content exceeds 0.3000%, seizure hardening is excessively developed as the amount of solid solution carbon in the steel increases, and the ductility of the steel sheet after baking hardening is significantly deteriorated. Therefore, the C content is set to 0.3000% or less.

"Si: more than 0% and less than 3.000%"
Si is an element that thermodynamically stabilizes ferrite, and is a useful element that dissolves in ferrite to develop solid solution hardening. In addition, Si is an element that promotes the formation of martensite by inhibiting the formation of carbides. In order to obtain these effects, the Si content is set to more than 0%.
On the other hand, Si is also an element that enhances the hardenability of steel. Therefore, if the Si content is too high, the transformation from austenite to ferrite is delayed in the cooling process after hot rolling, and bainite and / or martensite is likely to be formed. In order to generate a desired amount of ferrite in the cooling process after hot rolling, the Si content is 3.000% or less.

"Mn: 0.050% or more and 3.600% or less"
Mn is an element that thermodynamically stabilizes austenite. Further, Mn is an element useful for reducing surface defects during hot rolling due to S. In order to obtain these effects, the Mn content is set to 0.050% or more.
On the other hand, since Mn enhances hardenability, if the Mn content is too large, the transformation from austenite to ferrite is delayed in the cooling process after hot rolling, and bainite and / or martensite is likely to be formed. The Mn content is 3.600% or less in order to suppress an increase in the area ratio of bainite and / or martensite to obtain a desired steel structure.

"P: More than 0% and less than 0.030%"
P is an element that easily segregates at grain boundaries. If the P content is too high, embrittlement may be promoted in the high-strength steel sheet and the workability may be deteriorated. Therefore, the P content is 0.030% or less.
On the other hand, P is also an element that enhances the strength of the steel sheet by hardening the solid solution. Further, if the P content is excessively reduced, the refining cost increases. Therefore, the P content may be more than 0% or 0.001% or more.

"S: More than 0% and less than 0.0200%"
S exists as an inclusion in the steel. If the S content exceeds 0.0200%, the amount of inclusions may increase, and the inclusions may become the starting point of fracture during processing of the steel sheet, resulting in a decrease in formability of the steel sheet. Therefore, while the S content is 0.0200% or less, if the S content is excessively reduced, the refining cost increases. Therefore, the S content may be more than 0% or 0.0001% or more.

"Al: 0.500% or more and 5.000% or less"
Like Si, Al is an element that thermodynamically stabilizes ferrite and also inhibits the formation of carbides in steel. Further, Al is an element that increases the strength of the steel sheet when it is contained at the same time as a predetermined amount of Ni. It is considered that the increase in strength of the steel sheet by simultaneously containing Al and Ni is due to the formation of the columnar precipitation phase (Ni-Al intermetallic compound) of Ni—Al and the solid solution cluster. Further, in order to obtain the effect, it is considered necessary to promote the ferrite transformation after hot rolling.
If the Al content is less than 0.500%, a sufficient amount of ferrite is not produced, and the above effect cannot be obtained. Therefore, the Al content is set to 0.500% or more.
On the other hand, if the Al content exceeds 5.000%, ferrite will be generated during hot rolling. As a result, the processed austenite and the processed ferrite are formed during hot rolling, so that a mixed structure consisting of the ferrite produced by the transformation of the processed austenite during cooling and the ferrite produced by the recovery and recrystallization of the processed ferrite is formed. Will be done. The mechanical properties of such mixed structures vary sharply with their formation rate. Therefore, when the steel structure of the steel sheet has a mixed structure as described above, industrially stable mechanical properties cannot be obtained. Therefore, the Al content is set to 5.000% or less.

"N: More than 0% and less than 0.0100%"
N is an penetrating element and is an element that contributes to strain age hardening. However, since the high-strength steel sheet according to this embodiment contains Al as an essential element, N is mainly present as aluminum nitride (AlN) in the steel. When the N content exceeds 0.0100%, the coarse AlN content increases, and the coarse AlN may become the starting point of fracture during the processing of the steel sheet, and the formability of the steel sheet may deteriorate. Therefore, the N content is set to 0.0100% or less.
On the other hand, if the N content is excessively reduced, the refining cost increases. Therefore, the N content may be more than 0% or 0.0001% or more.

"Ni: 1.000% or more and 12.300% or less"
Like Mn, Ni is an element that thermodynamically stabilizes austenite and is also an element that inhibits the formation of carbides in steel. Further, Ni is an element that increases the strength of the steel sheet when it is contained at the same time as a predetermined amount of Al. The reason for this is not clear, but it is presumed that a columnar precipitation phase (Ni-Al intermetallic compound) of Ni and Al and a solid solution cluster are formed. The Ni content is 1.000% or more in order to obtain the effect of increasing the strength of the steel sheet by simultaneously containing Ni and Al.
On the other hand, when the Ni content exceeds 12.300%, the transformation from austenite to ferrite is delayed in the cooling process after hot rolling, and a large amount of bainite and / or martensite is formed to obtain a desired steel structure. do not have. Therefore, the Ni content is set to 12.300% or less.

In the composition of the high-strength steel sheet according to the present embodiment, elements other than the above-mentioned elements, that is, the balance is Fe and impurities. Impurities are elements that are inevitably mixed from steel raw materials or scrap, or elements that are inevitably mixed in the steelmaking process, and the high-strength steel sheet according to the present embodiment has the effect of the high-strength steel sheet according to the present embodiment. Elements that are acceptable within the playable range can be exemplified. The total content of impurities is preferably 0.100% or less. Of the impurities, the content of O is preferably 0.010% or less.
The high-strength steel sheet according to the present embodiment may contain one or more of the optional elements shown below instead of a part of the remaining Fe. However, since the high-strength steel sheet according to the present embodiment can solve the problem without containing any element shown below, the lower limit of the content of the arbitrary element is 0%.

"Cu: 0% or more and 4.800% or less"
Cu, like Mn and Ni, is an element that thermodynamically stabilizes the austenite phase. In order to surely obtain this effect, it is preferable that the Cu content is 0.050% or more.
On the other hand, if the Cu content is too high, a desired amount of ferrite cannot be produced in the cooling process after hot rolling. Therefore, even when it is contained, the Cu content is set to 4.800% or less.

"Mo: 0% or more and 2.500% or less"
Mo, like Si and Al, is an element that stabilizes the ferrite phase. In order to surely obtain this effect, it is preferable that the Mo content is 0.030% or more.
On the other hand, if the Mo content is too high, a desired amount of ferrite cannot be produced in the cooling process after hot rolling. Therefore, even when it is contained, the Mo content is 2.500% or less.

"Ca: 0% or more and 0.0200% or less"
"Mg: 0% or more and 0.0200% or less"
"REM: 0% or more and 0.0200% or less"
Ca, Mg and REM are all elements that control the shape of inclusions such as oxides and sulfides. Specifically, it is an element that finely disperses inclusions, reduces the factor of the starting point of fracture during processing, and contributes to the improvement of workability of the steel sheet. In order to surely exhibit these effects, it is preferable that the content of even one of Ca, Mg and REM is 0.0001% or more. However, if the content of even one of Ca, Mg and REM exceeds 0.0200%, the total number of inclusions increases and the internal quality of the steel sheet deteriorates. Therefore, the contents of Ca, Mg and REM are 0.0001% or more and 0.0200% or less, respectively.
REM refers to 17 elements from scandium, yttrium, lanthanum to lutetium, and is generally called a rare earth element or rare earth. When a plurality of types of REM are contained, the above-mentioned content of REM means the total content of these elements.

Among the above-mentioned elements, the contents of C, Si, Mn, Al, Ni, Cu and Mo need to satisfy the following equations (1), (2) and (3) at the same time.
(Ni% + 12.00 x C% + 2.00 x Mn% + 1.20 x Cu%) -2.00 x (Al% + 0.50 x Si% + 0.25 x Mo%) ≧ 0.0 (1)
3.725 x C% + 0.160 x Si% + 0.630 x Mn% -0.110 x Al% + 0.210 x Ni% + 0.450 x Cu% + 0.620 x Mo% -1.818 ≦ 1. 000
(2)
Ni% -0.5 x Cu% ≧ 0.0 (3)
Here, C%, Si%, Mn%, Al%, Ni%, Cu% and Mo% are the contents of C, Si, Mn, Al, Ni, Cu and Mo in mass%, respectively, and are contained. If not, substitute 0.

The above equation (1) is a conditional equation for hot rolling in the single-phase region of austenite without forming ferrite during hot rolling. Further, the formula (1) is also a conditional formula showing an index for improving the strength-ductility balance of the steel sheet by the columnar precipitation phase of Ni—Al and suppressing the decrease in ductility after baking and curing.
First, the former effect on the steel sheet characteristics will be described. If the content of the element that stabilizes austenite such as Si, Al and Mo is higher than the content of the element that stabilizes austenite such as C, Mn, Ni and Cu, not only austenite but also ferrite will be generated during hot rolling. Generated. As mentioned in the reason for limiting the Al content, austenite and ferrite mixed during hot rolling become processed austenite and processed ferrite by hot rolling, respectively. It becomes ferrite produced by crystals and forms a mixed structure. Since the mechanical properties of such mixed structures vary greatly depending on their formation ratio, industrially stable mechanical properties cannot be obtained.
Next, the latter effect on the steel sheet characteristics will be described. In the above-mentioned hot rolling, the Ni atom that stabilizes austenite is uniformly dispersed in the austenite phase, but the Al atom that stabilizes ferrite segregates at the grain boundaries of austenite rather than solid solution in the austenite phase. In the high-strength steel sheet according to the present embodiment, a columnar precipitation phase is generated at the ferrite grain boundary (including the vicinity) in the steel after hot rolling, so that the strength-ductility balance and the decrease in ductility after baking hardening are suppressed. To. Although the detailed mechanism is not clear, when a row-like precipitation phase is formed at the ferrite grain boundaries, the deformation resistance near the ferrite grain boundaries increases, and when the steel plate before the seizure hardening process is deformed, it is inside the ferrite grains of the steel. It is considered that the dislocations that proliferate are less likely to accumulate at the ferrite grain boundaries, so that the formation of voids due to deformation is suppressed and the strength-ductility balance is improved. Further, the cause of the deterioration of ductility after the baking hardening treatment in the conventional steel is that the dislocation mobility is lowered by applying the heat treatment for dislocation in a state where the dislocations are excessively accumulated at the ferrite grain boundaries, and the ferrite grain boundaries are lowered. It is understood that it is possible to suppress the decrease in ductility after the baking hardening treatment by suppressing the excessive accumulation of dislocations to. Therefore, in the high-strength steel sheet according to the present embodiment, in addition to the content of each element being in a predetermined range in the component composition, it is necessary to satisfy the above formula (1).

The above equation (2) is a conditional equation for generating ferrite during cooling after hot rolling. The present inventors have shown that the development of high strength by simultaneously containing Ni and Al is caused by the formation of a columnar precipitation phase at the ferrite grain boundaries after the formation of a ferrite structure having a bcc crystal structure. Is getting. When the amount of elements that enhance the hardenability exceeds a certain amount, ferrite is not sufficiently generated during cooling after hot rolling, and the formation of a columnar precipitate phase at the ferrite grain boundaries is suppressed, so that the strength of the steel sheet is increased and the steel sheet is strengthened. High ductility cannot be achieved. Therefore, in the component composition of the high-strength steel sheet according to the present embodiment, in addition to the content of each element being in a predetermined range, it is necessary to satisfy the above equation (2).

The above equation (3) is a conditional equation that defines the required Ni content from the viewpoint of preventing hot brittleness due to the molten phase of Cu. From the viewpoint of preventing hot brittleness by Cu, the Cu content is allowed up to 2.0 times the Ni content. When the Ni content is half (0.5 times) or more of the Cu content, hot brittleness can be suppressed. Therefore, it is necessary to satisfy the above equation (3) in the composition of the high-strength steel sheet according to the present embodiment.

Next, the steel structure (microstructure) of the high-strength steel sheet according to this embodiment will be described.

(Steel structure)
The steel structure of the high-strength steel plate according to the present embodiment has a total area ratio of ferrite, cementite and pearlite of 85.0% or more and 99.0% or less, and a total of 1.0% or more and 15.0% or less. Consists of martensite, bainite and retained austenite.

The area ratio of each tissue can be measured by the following method.
Samples are taken so that the cross section of the plate thickness parallel to the rolling direction of the steel plate is the observation surface. A region of 100 μm × 100 μm within the range of t / 8 to 3 t / 8 in the plate thickness direction centered on the position of 1/4 of the plate thickness t (position of t / 4) from the surface of the steel plate on the observation surface. The observation area. This observation area is corroded by repeller etching, and the tissue image of the corroded surface is analyzed using an optical microscope to calculate the area ratio of each tissue.

As described above, the steel structure of the high-strength steel plate according to the present embodiment is composed of a combination of a plurality of structures among ferrite, cementite, pearlite, bainite, martensite and retained austenite. The captured tissue image can be distinguished into a black part, a gray part, and a white part. Of these, the gray part contains ferrite and bainite, the black part contains cementite and pearlite, and the white part contains martensite and retained austenite. Therefore, the white part is regarded as martensite and retained austenite. In the gray part, the lath-like structure is regarded as bainite, and the structure other than the lath-like structure is regarded as ferrite. Further, the tissue arranged in dots in the black portion is regarded as cementite, and the tissue in which the black portions are arranged in rows or exist as a mass having a diameter of more than 1.0 μm is regarded as pearlite. The area ratio of each tissue is obtained by analyzing the tissue images of at least three observation regions and calculating the average value of the area ratio of each tissue.

"Area ratio of ferrite, cementite and pearlite: 85.0% or more and 99.0% or less in total"
From various experiments, the present inventors have found that the effect of increasing the strength during cooling is expressed by a ferrite or pearlite structure. That is, in order to sufficiently increase the strength of the steel sheet, it is important to include a desired amount of ferrite or pearlite. When ferrite and / or pearlite are produced in the composition of the high-strength steel plate according to the present embodiment, a small amount of cementite is also produced. Therefore, in this embodiment, the total area ratio of ferrite, cementite and pearlite is specified. In the present embodiment, the area ratios of ferrite, cementite and pearlite are 85.0% or more and 99.0% or less in total. In order to obtain a sufficient effect, the area ratio of ferrite alone in the above structure is preferably 40.0% or more.

"Area ratio of martensite, bainite and retained austenite: 1.0% or more and 15.0% or less in total"
In the steel structure of the high-strength steel plate according to the present embodiment, the structure other than ferrite, cementite and pearlite is one or more of martensite, bainite and retained austenite.
In this embodiment, it is important to include the desired amount of martensite, bainite and / or retained austenite. Therefore, in this embodiment, the total area ratio of martensite, bainite and retained austenite is defined. Yield elongation is significantly higher in steel structures consisting of ferrite, cementite and / or pearlite, free of martensite, bainite and retained austenite. In the portion where the yield elongation occurs, the deformed portion and the undeformed portion coexist, and the deformation is locally concentrated in the thin portion where the deformation has progressed, and the plate breakage occurs at an early stage. That is, the ductility of the steel sheet is low. Therefore, by adding 1.0% or more of martensite, bainite and retained austenite, it is possible to suppress an excessive increase in yield elongation. On the other hand, these structures contribute to the increase in strength, but as the area ratio increases, the amount of baking hardening also increases, and the ductility of the steel sheet after the coating baking treatment deteriorates. Therefore, the area ratio of martensite, bainite and retained austenite shall be 15.0% or less.

A row-like precipitation phase of Ni—Al exists on the grain boundaries of ferrite, and the coverage of the ferrite grain boundaries by the row-like precipitation phase is 5.0% or more of the total grain boundary length. Is present on the grain boundaries of ferrite (covers the ferrite grain boundaries), and the coverage of the ferrite grain boundaries by the columnar precipitation phase of Ni—Al is 5.0% or more, which is high according to the present embodiment. With a strong steel plate, high strength of 590 MPa or more can be obtained even if the area ratio of structures such as baynite and martensite is small.
The upper limit of the coverage is not particularly limited, but is preferably 90.0% or less in terms of suppressing an excessive increase in yield elongation.

The identification of the columnar precipitation phase of Ni—Al and the measurement of the coverage of the ferrite grain boundaries are performed by the microstructure observation photograph obtained by a transmission electron microscope (TEM) as shown below. In the present embodiment, Ni-Al intermetallic compounds having a rod-like shape having a diameter of 20 nm or less and a length of 500 nm or less are precipitated in a row or a dot shape, and the distance between adjacent Ni-Al intermetallic compounds. The massive structure dispersed so as to be within 100 nm is referred to as one columnar precipitation phase.
First, a material for observation is cut out from the steel plate using a precision cutting machine, and this is cut and polished with emery paper to a thickness of 1/4 of the plate thickness direction, which is the observation position, and the material thickness is adjusted to 0.1 mm. After that, a sample for TEM observation is prepared by performing double-sided jet electrolytic polishing on a sample punched from a material having a diameter of 3 mm by punching. In this TEM observation sample, a region of 4.0 μm × 4.0 μm was randomly selected, and Ni-Al was analyzed by elemental analysis by EDS and crystal structure analysis by nanobeam diffraction (NBD / Nano Beam Diffraction) method. The location of the formation of the intermetallic compound in the above is identified, and a TEM photograph is taken so that the location of the formation is included. In 25 TEM photographs obtained by repeating this operation, the grain boundary length of ferrite in contact with the columnar precipitation phase of Ni—Al is divided by the total grain boundary length of ferrite to cover the ferrite grain boundary. Find the rate.
The length of the ferrite grain boundaries in contact with the columnar precipitation phase is determined by the following procedure by image analysis. In the columnar precipitation phase (region where the intermetallic compound is dispersed and precipitated) existing in one ferrite grain, a circle with a radius of 100 nm is drawn around each of the Ni—Al intermetallic compounds located on the outermost periphery thereof. , The range obtained by connecting these circles (the range included in any of the circles) is defined as the boundary region of the columnar precipitation phase, and the length of the ferrite grain boundary contained in this boundary region is defined as the boundary region of the columnar precipitation phase. The length of the ferrite grain boundary in contact with.
The ferrite grain boundary is a boundary in a ferrite phase having a body-centered cubic lattice crystal structure with a crystal orientation difference of 15 ° or more. It can be obtained from the identification of the orientation.

(Plating layer)
The high-strength steel sheet according to the present embodiment may be provided with a plating layer on its surface for the purpose of improving corrosion resistance and the like. That is, the high-strength steel sheet according to the present embodiment may include a hot-dip galvanized layer, an alloyed hot-dip galvanized layer, or an electrozinc plated layer. The amount of adhesion of the plating layer is not particularly limited, and may be a general amount of adhesion.

(Production method)
Next, a preferable manufacturing method for obtaining a high-strength steel sheet according to the present embodiment will be described. The following description is intended merely as an example of a manufacturing method for obtaining a high-strength steel sheet according to the present embodiment, and is not intended to be limited to the following manufacturing method.

Conventionally, in order to exert the effect of increasing the strength by the intermetallic compound of Ni-Al, a steel sheet whose temperature has dropped to room temperature, as typified by Maruage steel, is used in a temperature range of, for example, 450 to 550 ° C. It was necessary to apply aging treatment for about 5 hours. This method requires an additional step of applying heat treatment in the production of hot-rolled steel sheet, and the productivity is low. The present inventors have found that by controlling the composition and the cooling history after hot rolling within an appropriate range, the strength of the steel sheet can be increased without performing additional heat treatment. It was also found that, in the conventional method, the intermetallic compound is not always formed as a columnar precipitation phase so as to cover the grain boundaries, and a sufficient strength improving effect may not be obtained. The method for manufacturing a high-strength steel sheet according to the present embodiment obtained based on this finding will be described in detail below.
The temperature described below refers to the surface temperature of the steel sheet.

In the method for producing a high-strength steel sheet according to the present embodiment, first, a steel piece having the above-mentioned composition is produced. The method for producing this piece of steel is not particularly limited. For example, the steel pieces may be manufactured by a general method such as continuous casting or thin slab casters.

The obtained steel piece is used as it is, or once the temperature of the steel sheet is lowered to room temperature, the steel piece is heated. The heating temperature at this time is preferably 900 ° C. or higher and 1300 ° C. or lower. If the heating temperature is less than 900 ° C., the austenite transformation cannot be sufficiently performed, the solution of the precipitate becomes insufficient, and the desired strength may not be obtained. On the other hand, when the heating temperature exceeds 1300 ° C., the amount of scale generated increases, and the surface defects associated therewith may increase. The heating time in the temperature range of 900 ° C. or higher and 1300 ° C. or lower is preferably 60 minutes or longer in order to sufficiently transform austenite. Further, the heating time is preferably 240 minutes or less from the viewpoint of suppressing scale formation.

After heating the steel pieces, hot rolling is performed. In this hot rolling step, the final rolling temperature is set to 750 ° C. or higher. By terminating the hot rolling at a temperature of 750 ° C. or higher, it is possible to suppress the formation of ferrite during the hot rolling and the mixture of processed austenite and processed ferrite. When the processed austenite and the processed ferrite are cooled later, the former becomes ferrite produced by transformation and the latter becomes ferrite produced by recovery / recrystallization, forming a mixed structure. Since the mechanical properties of such a mixed structure vary greatly depending on the formation ratio, industrially stable mechanical properties cannot be obtained. Further, the processed ferrite is in a state where high-density dislocations are left in the ferrite grains, and the above-mentioned columnar precipitation phase of Ni—Al is preferentially generated on the dislocations in the ferrite grains rather than the ferrite grain boundaries. Therefore, when processed ferrite is generated, the formation of a columnar precipitation phase of Ni—Al at the ferrite grain boundaries is suppressed. Therefore, the final rolling temperature is set to 750 ° C. or higher. The final rolling temperature is the surface temperature of the steel sheet on the exit side of the finish rolling mill.

After hot rolling, the steel sheet is cooled. This cooling step includes a first cooling step of cooling so that the average cooling rate in a temperature range of 550 ° C. or higher and 750 ° C. or lower is 1 ° C./sec or higher and 20 ° C./sec or lower, and a temperature range of 550 ° C. or lower and 300 ° C. or higher. It consists of a second cooling step of cooling to 300 ° C. or lower so that the average cooling rate in the above is 30 ° C./sec or more and 300 ° C./sec or less. The average cooling rate in the first cooling step is a value obtained by dividing the temperature drop width from 750 ° C. to 550 ° C. by the time required for cooling from 750 ° C. to 550 ° C. The average cooling rate in the second cooling step is a value obtained by dividing the temperature drop width from 550 ° C to 300 ° C by the time required for cooling from 550 ° C to 300 ° C.
The average cooling rate in the first cooling step and the second cooling step is adjusted, for example, by adjusting the amount of water jetted from the nozzle toward the steel plate on the cooling floor.

In the first cooling step, cooling is performed so that the average cooling rate in the temperature range of 550 ° C. or higher and 750 ° C. or lower is 1 ° C./sec or more and 20 ° C./sec or lower. By setting the average cooling rate in the temperature range of 550 ° C or higher and 750 ° C or lower to 20 ° C / sec or lower, ferrite is generated from the inside of the austenite grains, and Ni is covered with the ferrite grain boundaries in the winding step described later. -The columnar precipitation phase of Al is easily formed, and the strength of the steel plate can be increased. The slower the cooling rate in the above temperature range, the higher the strength of the steel sheet. However, if the cooling rate in the above temperature range is too slow, the amount of scale generated on the surface of the steel sheet increases, and the cost in the pickling process for removing surface defects and scale increases, which is not preferable in terms of production cost. .. Therefore, the average cooling rate in the above temperature range is preferably 1 ° C./sec or more.

In the second cooling step, the steel sheet is cooled to 300 ° C. or lower so that the average cooling rate in the temperature range of 300 ° C. or higher and 550 ° C. or lower is 30 ° C./sec or more and 300 ° C./sec or lower. Then, the steel sheet is wound at 300 ° C. or lower. By setting the average cooling rate in the temperature range of 300 ° C. or higher and 550 ° C. or lower to 30 ° C./sec or higher, austenite that has not been transformed up to 550 ° C. is transformed into bainite and / or martensite. If the area ratio of bainite and / or martensite at this time is large, the amount of baking hardening increases, the strength of the steel sheet increases remarkably, and the ductility after baking hardening decreases. On the other hand, in the structure which does not contain martensite and bainite and is composed of ferrite, cementite and pearlite, yield elongation is likely to occur. In the portion where the yield elongation occurs, a portion where the deformation has progressed and a portion where the deformation is small occurs (non-uniform deformation occurs), and a difference in plate thickness occurs in the steel sheet. If it is further deformed in this state, work hardening starts, but the ductility deteriorates because the deformation progresses concentrated on the portion where the plate thickness is thin. Therefore, it is necessary to include a desired amount of bainite and / or martensite to suppress an excessive increase in yield elongation and promote uniform deformation. In order to obtain the desired amount of bainite and / or martensite, the average cooling rate in the temperature range of 300 ° C. or higher and 550 ° C. or lower is set to 30 ° C./sec or higher. When the average cooling rate in the above temperature range exceeds 300 ° C., cooling unevenness is likely to occur and the shape of the steel sheet may be deteriorated. Therefore, the average cooling rate in the above temperature range is preferably 300 ° C./sec or less.

After the second cooling step, the steel sheet is wound up. The winding temperature of the steel sheet is preferably 50 ° C. or higher and 300 ° C. or lower. The reason for this is that fine Ni-Al columnar precipitation phases are generated at the ferrite grain boundaries, the strength-ductility balance of the hot-rolled steel sheet is improved, and the effect of suppressing deterioration of ductility after baking and curing is remarkably obtained. Because it can be done. Although the grain size of the ferrite grains generated in the austenite grains in the above-mentioned first cooling step increases during the first cooling step, the movement of the austenite / ferrite interface at this time segregates Al atoms and the like. Pinned at the austenite grain boundaries. Further, since ferrite formation from the inside of the grain and pinning of the austenite / ferrite interface movement at the austenite grain boundary occur in each of the adjacent austenite grains, the adjacent austenite grains are present at the austenite grain boundary where Al atoms and the like are segregated. The ferrite interface generated from the inside comes into contact with each other to form ferrite grain boundaries. Since Al is concentrated at the ferrite grain boundaries thus generated, by winding at 50 ° C. or higher and 300 ° C. or lower, the Ni-Al columnar precipitation phase using the Al concentrated at the ferrite grain boundaries is used. Is generated starting from the ferrite grain boundary. When the winding temperature exceeds 300 ° C., the time for holding at a high temperature increases, so that Al segregated at the ferrite grain boundaries diffuses from the grain boundaries into the grains, and the Al concentration at the ferrite grain boundaries decreases. By doing so, it is considered that the coverage of ferrite by the columnar precipitation phase of Ni—Al decreases. Further, when the winding temperature is less than 50 ° C., the diffusion frequency of Al concentrated in the ferrite grain boundaries into the ferrite grains is significantly reduced, so that the coverage of the ferrite grain boundaries by the columnar precipitation phase of Ni—Al is reduced. Is thought to decrease.

When a hot-dip galvanized steel sheet is obtained by forming a hot-dip galvanized steel sheet on the high-strength steel sheet according to the present embodiment, the hot-dip galvanized steel sheet is unwound as described above, then descaled, and then hot-dip galvanized. Should be applied. When a hot-dip galvanized steel sheet is obtained by forming an alloyed hot-dip galvanized layer, the hot-dip galvanized steel sheet may be alloyed. The thermal stability of the columnar precipitation phase of Ni—Al produced by hot rolling is high, and the columnar precipitation phase formed at the ferrite grain boundaries is not completely melted unless it is heated to 1000 ° C. or higher. When heated to a temperature of 1000 ° C. or higher, the austenite grains become abnormally coarse, and even if the steel sheet is subsequently given a cooling history within the range shown by the hot rolling conditions of the present invention, the structural morphology shown in the present invention cannot be obtained. Therefore, in order to obtain the structure of the present invention even after the plating treatment, when the steel sheet before the plating treatment is heated, the temperature is limited to the temperature range of 400 ° C. or higher and 730 ° C. or lower, which is the ferrite region. There is a need to.

Next, the effect of one aspect of the present invention will be described more specifically by way of examples, but the conditions in the examples are one condition example adopted for confirming the feasibility and effect of the present invention. The present invention is not limited to this one-condition example. The present invention may adopt various conditions as long as the gist of the present invention is not deviated and the object of the present invention is achieved.

(Example 1)
Steels having the component compositions shown in Tables 1A and 1B were melted, cast, and solidified to obtain steel pieces. Then, the steel pieces are hot-rolled as they are, or once cooled to room temperature, heated at 1050 ° C or higher and 1300 ° C or lower for 60 minutes or longer and 240 minutes or shorter, and hot-rolled under the conditions shown in Tables 2A and 2B. Manufactured steel plate. The thickness of the hot-rolled steel sheet was 3.2 mm. The values on the left side of each of the above equations (1), (2) and (3) are also shown in Tables 1A and 1B.

Yield strength (YP) and tensile (maximum) strength (TS) are obtained by collecting test pieces from the manufactured hot-rolled steel sheet so that the direction perpendicular to the rolling direction is the longitudinal direction and conducting a tensile test. , Ell was measured. Here, the yield strength, tensile strength and elongation were measured using the No. 5 test piece described in JIS Z 2241: 2011 according to the method described in JIS Z 2241: 2011. When the tensile strength (TS) was 590 MPa or more, it was judged to have a desired strength and was judged to be acceptable, and when it was less than 590 MPa, it was judged to have no desired strength and was judged to be unacceptable. TS × El (MPa ·%) was determined from the obtained tensile strength (TS) and elongation (El), and the strength-ductility balance was evaluated. When TS × El was 15,000 MPa ·% or more, it was judged that the strength-ductility balance was excellent.
Further, after applying 2% tensile deformation to the No. 5 test piece collected from the hot-rolled steel sheet so that the direction perpendicular to the rolling direction is the longitudinal direction, strain aging hardening treatment at 170 ° C. for 20 minutes (hereinafter referred to as “1”). BH treatment) was applied. This strain age hardening treatment was performed to simulate the coating baking treatment. Then, by carrying out a tensile test according to the method described in JIS Z 2241: 2011, the tensile strength (TS), yield strength (YP) and elongation (El) after the BH treatment were measured. When the tensile strength (TS) after the BH treatment was 590 MPa or more, it was judged to be acceptable because it had the desired strength, and when it was less than 590 MPa, it was judged to be unacceptable because it did not have the desired strength. When the reduction margin of elongation before and after the BH treatment was 4% or less, it was judged that the decrease in ductility due to baking hardening was suppressed, and it was judged to be acceptable. On the other hand, when the reduction margin of elongation was more than 4%, it was judged that the decrease in ductility due to baking hardening was not suppressed, and it was judged to be unacceptable.
The area ratio of each structure in the hot-rolled steel sheet was obtained by the above-mentioned method.
The above results are shown in Tables 2A and 2B.

No. of the present invention example. It can be seen that 1 to 14 have a tensile strength of 590 MPa or more, and even if the yield strength increases after the BH treatment, the deterioration of ductility is small. Further, since the strength-ductility balance is excellent, it is considered that the press formability is also excellent.
No. Reference numerals 15 to 36 are comparative examples in which one or more of the strength and ΔEl are inferior in characteristics.

No. In No. 15, since the C content was low, sufficient amounts of martensite, bainite and retained austenite could not be obtained, and the ductility after BH treatment deteriorated.
No. In No. 16, since the C content was high, the BH amount after the BH treatment increased, the yield strength and the tensile strength increased remarkably, and the ductility deterioration was large.
No. 17 and No. No. 18 had a large Si content and Mn content, respectively. Since 18 did not satisfy the above equation (2), the area ratio of martensite was large, the strength increase allowance after the BH treatment was large, and the ductility deterioration was large.
No. 19 and No. Since 20 had a large P content and an S content, respectively, the ductility deterioration was large.
No. 21, No. 24 and No. Since the Al content and the Ni content of each of 33 were low, the tensile strength before and after the BH treatment was less than 590 MPa, resulting in insufficient strength. Further, since the coverage of the ferrite grain boundaries due to the columnar precipitation phase of Ni—Al decreased, the ductility deterioration was large.
No. Since the Al content of 22 was high, ferrite was generated during hot rolling, so that the processed ferrite processed by rolling and the ferrite generated in the cooling process after hot rolling had a structure after winding. Although the strength increased due to the inclusion of Al due to the non-uniform mixing, the ductility was significantly inferior due to the non-uniformity of the structure, and the ductility was greatly deteriorated.
No. Since 23 had a high N content, coarse aluminum nitride (AlN) was generated and became a starting point of fracture during processing, the strength-ductility balance was significantly inferior, and the ductility deterioration was large.
No. 25, No. 26, No. 27 and No. Since each of 34 had a large Ni content, Cu content and Mo content and did not satisfy the above formula (2), martensite was generated, the strength increase margin after the BH treatment was large, and the ductility deterioration was large.
No. Since 28 did not satisfy the above equation (1), ferrite was generated during hot rolling, and the structure became non-uniform including the recovered / recrystallized ferrite, so that the ductility was inferior and the ductility was greatly deteriorated.
No. In No. 29, the composition of the steel was in a preferable range, but since the final rolling temperature was low, ferrite was generated during the hot rolling, and the steel was mixed unevenly in the structure after winding, so that the strength was increased. It was insufficient and the ductility deteriorated.
No. In No. 30, the composition of the steel was in a preferable range, but the average cooling rate at 750 to 550 ° C. was high, so that the amount of martensite increased and the ductility after the BH treatment was greatly deteriorated.
No. In No. 31, the composition of the steel was in a preferable range, but the average cooling rate at 550 to 300 ° C. was small, so that the desired amounts of martensite, bainite and retained austenite could not be obtained, and the ductility deterioration was large.
No. In No. 32, the composition of the steel was in a preferable range, but the winding temperature was low, the proportion of ferrite in the steel structure decreased, and at the same time, the coverage of ferrite grain boundaries due to the columnar precipitation phase of Ni—Al decreased. Therefore, the ductility deterioration was large.
No. In No. 35, since the above equation (2) was not satisfied, martensite was generated, the strength increase allowance after the BH treatment was large, and the ductility deterioration was large.
No. In No. 36, the composition of the steel was in a preferable range, but the winding temperature was high and the coverage of the ferrite grain boundaries due to the columnar precipitation phase of Ni—Al was lowered, so that the ductility deterioration was large.

(Example 2)
The hot-rolled steel sheets of sample codes 1 and 8 produced in Example 1 were heated to 660 to 720 ° C. and subjected to hot-dip galvanizing treatment to obtain hot-dip galvanized steel sheets. Sample code 8 was formed into an alloyed hot-dip galvanized steel sheet by performing an alloying treatment at 540 to 580 ° C. after the hot-dip galvanizing treatment. The obtained hot-dip galvanized steel sheet (sample code 1) and alloyed hot-dip galvanized steel sheet (sample code 8) were subjected to microstructure observation and mechanical property evaluation in the same manner as in Example 1. The results are shown in Table 3.

As shown in Table 3, the same results as in the examples were obtained for any of the plated steel sheets, and it was confirmed that the desired characteristics in the present invention can be ensured even if the hot dip galvanizing treatment or the alloying treatment is performed. did it.

According to the above aspect according to the present invention, it is possible to provide a high-strength steel sheet in which a decrease in ductility due to baking hardening is suppressed. The high-strength steel sheet according to the above aspect is suitable for a shock absorbing member that requires collision resistance. For example, by applying it to parts such as automobiles and transportation equipment, fuel efficiency is improved by reducing the weight of the vehicle body and collision safety is achieved. Since it can contribute to the further improvement of the industrial value, it has extremely high industrial utility value.

Claims (3)

Ingredient composition is mass%,
C: 0.0150% or more and 0.3000% or less,
Si: More than 0% and less than 3.000%,
Mn: 0.050% or more and 3.600% or less,
P: More than 0% and 0.030% or less,
S: More than 0% and 0.0200% or less,
Al: 0.500% or more and 5.000% or less,
N: More than 0% and less than 0.0100%,
Ni: 1.000% or more and 12.300% or less,
Cu: 0% or more and 4.800% or less,
Mo: 0% or more and 2.500% or less,
Ca: 0% or more and 0.0200% or less,
Mg: 0% or more and 0.0200% or less, and REM: 0% or more and 0.0200% or less, and the balance consists of Fe and impurities.
Satisfying the relationships of the following equations (1), (2) and (3) at the same time,
Steel structure from ferrite, cementite and pearlite totaling 85.0% or more and 99.0% or less, and martensite, bainite and retained austenite totaling 1.0% or more and 15.0% or less in area ratio. Configured,
The columnar precipitation phase of Ni—Al is present on the grain boundaries of ferrite, and the coverage of the ferrite grain boundaries by the columnar precipitation phase is 5.0% or more of the total grain boundary length.
A high-strength steel sheet having a tensile strength of 590 MPa or more.
(Ni% + 12.00 x C% + 2.00 x Mn% + 1.20 x Cu%) -2.00 x (Al% + 0.50 x Si% + 0.25 x Mo%) ≧ 0.0 (1)
3.725 x C% + 0.160 x Si% + 0.630 x Mn% -0.110 x Al% + 0.210 x Ni% + 0.450 x Cu% + 0.620 x Mo% -1.818 ≦ 1. 000
(2)
Ni% -0.5 x Cu% ≧ 0.0 (3)
Here, C%, Si%, Mn%, Al%, Ni%, Cu% and Mo% are the contents of C, Si, Mn, Al, Ni, Cu and Mo in mass%, respectively, and are contained. If not, substitute 0. The composition of the components is mass%.
Cu: 0.050% or more and 4.800% or less,
Mo: 0.030% or more and 2.500% or less,
Ca: 0.0001% or more and 0.0200% or less,
The high strength according to claim 1, wherein Mg: 0.0001% or more and 0.0200% or less, and REM: 0.0001% or more and 0.0200% or less are contained at least one selected. Steel plate. The high-strength steel sheet according to claim 1 or 2, which has a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface.