KR20240035536A - High-strength steel sheet, high-strength plated steel sheet and their manufacturing method, and members - Google Patents

Abstract

The present invention provides a high-strength steel sheet with a high yield ratio that has a tensile strength of 1180 MPa or more and has excellent bendability and toughness, and a method for manufacturing the same. The high-strength steel sheet of the present invention contains C, Si, Mn, P, S, Al, N, Ti, Nb, and B, the balance is Fe and inevitable impurities, and the components satisfy the following formula (1) It has a composition, the total area ratio of martensite and bainite is 95% or more, the average grain size of prior austenite grains is 10 μm or less, the B concentration of prior austenite grain boundaries is 0.10% or more in mass%, and martenite It has a C-enriched region along the site grain boundary, and the C-enriched region has a C concentration of 4.0 times or more than the C content in the steel, and has an enrichment width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary. , and also has a length of 100 nm or more in the direction parallel to the martensite grain boundary. ([％N]/14)/([％Ti]/47.9)＜1.0… (One)

Description

High-strength steel sheet, high-strength plated steel sheet and their manufacturing method, and members

The present invention relates to high-strength steel sheets, their manufacturing methods, and members.

In steel sheets for automobiles, increased strength is required to improve fuel efficiency by reducing the weight of the vehicle body. For skeleton parts, high-strength steel sheets with a tensile strength of 1180 MPa or more are required. In addition, in order to press-process a steel sheet into a desired shape, the steel sheet is required to have high bendability. Additionally, from the viewpoint of crash safety, in addition to strength, there are automobile parts that are required not to be easily deformed so as to secure living space for occupants in the event of a crash. For these automobile parts, the use of steel sheets with a high yield ratio is desired. In addition, high toughness is required so that automobile parts do not fracture in the event of a collision.

Patent Document 1 discloses a high-strength steel plate with excellent processability and low-temperature toughness and a method for manufacturing the same. Patent Document 2 discloses a high-strength steel sheet with excellent formability and impact resistance, and a method for manufacturing a high-strength steel sheet with excellent formability and impact resistance. Patent Document 3 discloses a high-yield ratio high-strength steel plate and its manufacturing method.

Japanese Patent No. 5728108 Specification Japanese Patent No. 6597939 Specification Japanese Patent No. 6700398 Specification

However, in Patent Documents 1 and 2, the yield ratio is not considered. In Patent Document 3, toughness is not considered.

As described above, it is difficult to manufacture a high-strength steel sheet with a tensile strength of 1180 MPa or more, excellent bendability and toughness, and a high yield ratio using the prior art.

The present invention was made in consideration of these circumstances, and its purpose is to provide a high-yield steel sheet with a tensile strength of 1180 MPa or more, which has excellent bendability and toughness, and a method for manufacturing the same.

In addition, in the present invention, high strength means that the tensile strength TS measured based on JIS Z2241 is 1180 MPa or more.

In addition, excellent bendability means that there is no cracking at the ridge portion of the bending peak in a bending test conducted in accordance with JISZ2248.

In addition, excellent toughness means that the brittle-ductile transition temperature is -40°C or lower in a Charpy impact test conducted in accordance with JIS Z2242.

In addition, high yield ratio means that the ratio YS/TS of yield strength and tensile strength measured based on JIS Z 2241 is 0.80 or more.

In order to achieve the above-described problem, the present inventors conducted intensive studies and obtained the following findings.

(1) The occurrence and propagation of cracks during bending and the failure path during brittle fracture follow the old austenite grain boundaries. Therefore, it is effective to improve bendability by refining the crystal grain size, complicating the fracture path, and increasing the strength of the grain boundaries. To refine the old austenite grains, it is effective to set the annealing temperature as low as possible above 850°C, which is the austenite single phase range. On the one hand, grain boundary segregation of B is effective in strengthening grain boundaries, but the amount of grain boundary segregation of B increases as annealing is performed at a higher temperature. Therefore, in order to increase the amount of grain boundary segregation of B while maintaining a fine crystal grain size, fine austenite grains are obtained by annealing at around 850°C, and then rapid heating and rapid cooling are performed. Accordingly, grain boundary segregation of B by diffusion is promoted while suppressing grain growth, and refinement of the austenite grain size and grain boundary segregation of B can be achieved simultaneously.

(2) The dislocations existing in the quenched martensite structure are mobile dislocations that easily cause sliding movement with low stress, and for this reason, the yield stress of the martensite structure is low. However, when the steel sheet after quenching is slightly processed, these dislocations move near the grain boundaries and become entangled there, forming immobile dislocations. Accordingly, the yield ratio of the steel plate can be increased.

(3) By tempering the steel sheet at low temperature, carbon segregates or clusters precipitate on the dislocations. By tempering a processed steel sheet in which dislocations have accumulated near grain boundaries at a low temperature, a region with high C concentration (C-enriched region) is formed on the network along the grain boundaries, significantly increasing the strength near the grain boundaries. Since C concentrates not only on the grain boundaries but also on the matrix phase between the grain boundaries, the effect of increasing strength is very large. Since the grain boundaries are not easily deformed, the yield ratio also increases significantly due to the formation of the C-enriched region.

The present invention has been made based on the above recognition. That is, the main structure of the present invention is as follows.

[1] In mass%,

C: 0.10% or more and 0.30 or less,

Si: 0.20% or more and 1.20% or less,

Mn: 2.5% or more and 4.0% or less,

P: 0.050% or less,

S: 0.020% or less,

Al: 0.10% or less,

N: 0.01% or less,

Ti: 0.100% or less,

Nb: 0.002% or more and 0.050% or less and

B: 0.0005% or more and 0.0050% or less

Contains, the balance consists of Fe and inevitable impurities, and has a component composition that satisfies the following formula (1),

The total area ratio of martensite and bainite is 95% or more,

The average grain size of the old austenite grains is 10㎛ or less,

The B concentration of the old austenite grain boundaries is 0.10% or more in mass%,

There is a C-enriched region along the martensite grain boundary,

The C concentration in the C-enriched area is 4.0 times or more than the C content in the steel,

A high-strength steel sheet having a thickening width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundaries, and having a length of 100 nm or more in the direction parallel to the martensite grain boundaries.

([％N]/14)/([％Ti]/47.9)＜1.0… (One)

In formula (1), [%N] and [%Ti] represent the content (% by mass) of N and Ti in the steel, respectively.

[2] The above component composition is further expressed in mass%,

V: 0.100 or less,

Mo: 0.500% or less,

Cr: 1.00% or less,

Cu: 1.00% or less,

Ni: 0.50% or less,

Sb: 0.200% or less,

Sn: 0.200% or less,

Ta: 0.200% or less,

W: 0.400% or less,

Zr: 0.0200% or less,

Ca: 0.0200% or less,

Mg: 0.0200% or less,

Co: 0.020% or less,

REM: 0.0200% or less,

Te: 0.020% or less,

Hf: 0.10% or less and

Bi: 0.200% or less

The high-strength steel sheet according to [1] above, which contains at least one element selected from the group consisting of:

[3] A high-strength plated steel sheet having a plating layer on at least one side of the high-strength steel sheet according to [1] or [2] above.

[4] A steel slab having the composition of [1] or [2] above is subjected to hot rolling to obtain a hot rolled sheet,

Cold rolling is performed on the hot-rolled sheet to obtain a cold-rolled sheet,

The cold-rolled sheet is heated to a first heating temperature of 850°C or more and 920°C or less and held for 10 s or more, and then heated to a second heating temperature of 1000°C or more and 1200°C or less at an average heating rate of 50°C/s or more. An annealing process is performed in which the temperature is raised to and cooled to 500°C or lower at an average cooling rate of 50°C/s or more within 5 seconds after reaching the second heating temperature,

After the annealing process, a rolling process is performed to obtain a second cold-rolled sheet by rolling the cold-rolled sheet to an elongation of 0.5% or more,

A method of producing a high-strength steel sheet, in which, after the rolling process, a reheating process is performed in which the second cold-rolled sheet is maintained at a reheating temperature of 70°C or more and 200°C or less for 600s or more to obtain a high-strength steel sheet.

[5] After the annealing process described in [4] above, but before the reheating process, plating is applied to at least one side of the high-strength steel sheet to obtain a high-strength plated steel sheet. A method for producing a high-strength plated steel sheet, including a plating process.

[6] A member formed by using at least part of the high-strength steel plate described in [1] or [2] above.

[7] A member formed by using at least part of the high-strength plated steel sheet described in [3] above.

According to the present invention, it is possible to provide a high-strength steel sheet with a high yield ratio that has a tensile strength of 1180 MPa or more and has excellent bendability and toughness, and a method for manufacturing the same.

1 is a diagram showing an example of a C-enriched region.

(Form for carrying out the invention)

Hereinafter, embodiments of the present invention will be described. Additionally, the present invention is not limited to the following embodiments. First, the appropriate range of the component composition of the steel sheet and the reason for its limitation will be explained. In addition, in the following description, “%” indicating the content of the component elements of the steel sheet means “mass%” unless otherwise specified. In addition, in this specification, the numerical range indicated using “~” means a range including the numerical values written before and after “~” as the lower limit and upper limit.

C: 0.10% or more and 0.30 or less

In addition to strengthening the martensite structure and bainite structure, C segregates at dislocations accumulated near the old austenite grain boundaries to strengthen the grain boundaries, and has the effect of increasing bendability, toughness, and yield ratio. If the C content is less than 0.10%, the area ratio of martensite and bainite decreases, and TS of 1180 MPa or more cannot be obtained. If the C content exceeds 0.30%, carbon boride of B and iron is formed during annealing, and a sufficient amount of B cannot be segregated at the old austenite grain boundaries. The C content is preferably 0.11% or more. Additionally, the C content is preferably 0.28% or less.

Si: 0.20% or more and 1.20% or less

Si is an element effective for solid solution strengthening, and requires addition of 0.20% or more. Meanwhile, Si is an element that stabilizes ferrite and increases its transformation point. Therefore, when the Si content exceeds 1.20%, it becomes difficult to reduce the prior austenite grain size to 10 μm or less. The Si content is preferably 0.50% or more. The Si content is preferably 1.10% or less.

Mn: 2.5% or more and 4.0% or less

Mn is effective in improving quenching properties. If the Mn content is less than 2.5%, the area ratio of martensite and bainite decreases, and the strength decreases. On the other hand, if the Mn content exceeds 4.0%, the segregated portion becomes excessively hard and the bendability decreases. The Mn content is preferably 2.8% or more. The Mn content is preferably 3.5% or less.

P: 0.050% or less

Since P segregates at the old austenite grain boundaries and reduces toughness, the P content is set to 0.050% or less. The lower limit of the P content is not particularly set, and may be 0%. However, setting it below 0.001% increases manufacturing costs, so 0.001% or more is preferable. The P content is preferably 0.025% or less.

S: 0.020% or less

Since S segregates at old austenite grain boundaries and reduces toughness, it is set to 0.020% or less. There is no particular lower limit for the S content, but setting it to less than 0.0001% increases manufacturing costs, so it is preferable to set it to 0.0001% or more. The S content is preferably 0.018% or less.

Al: 0.10% or less

Al is an element that acts as a deoxidizing agent, and in order to obtain such an effect, the Al content is preferably set to 0.005% or more. On the other hand, if the Al content exceeds 0.10%, ferrite is easily formed and the strength decreases. The Al content is preferably 0.05% or less.

N: 0.01% or less

N forms nitride with Nb and B, reducing the effect of addition of Nb and B. Therefore, it is set to 0.01% or less. There is no particular lower limit, but it is preferably 0.0001% or more from the viewpoint of manufacturing cost.

Ti: 0.100% or less

Ti has the effect of fixing N in the steel as TiN and suppressing the production of BN and NbN. In order to obtain these effects, the Ti content is preferably 0.005% or more. On the other hand, if the Ti content exceeds 0.100%, coarse Ti carbides are formed on the grain boundaries, and the toughness decreases. The Ti content is preferably 0.05% or less.

Nb: 0.002% or more and 0.050% or less

Nb precipitates as a solid solution or fine carbide and suppresses the growth of austenite grains during annealing. To obtain this effect, the Nb content is set to 0.002% or more. On the other hand, if the Nb content exceeds 0.050%, not only is the effect saturated, but coarse Nb carbides are precipitated, thereby reducing toughness. The Nb content is preferably 0.005% or more. Moreover, the Nb content is preferably 0.040% or less.

B: 0.0005% or more and 0.0050% or less

B has the effect of segregating at the old austenite grain boundaries and increasing the grain boundary strength. To achieve such effects, the B content is set to 0.0005% or more. On the other hand, if the B content exceeds 0.0050%, carbon boride is formed and toughness decreases. The B content is preferably 0.0010% or more. Additionally, the B content is preferably 0.0030% or less.

([％N]/14)/([％Ti]/47.9)＜1.0… (One)

In order to obtain the effect of adding B and Nb described above, N, which easily combines with these elements, needs to be fixed by Ti. Therefore, the mole fraction of N is made smaller than the mole fraction of Ti. That is, the contents of N and Ti in the steel are adjusted to satisfy the above equation (1). In addition, in Formula (1), [%N] and [%Ti] represent the content (% by mass) of N and Ti in the steel, respectively.

[Random Component]

In addition to the above component composition, the high-strength cold-rolled steel sheet according to the present embodiment further has, in mass%, V: 0.100% or less, Mo: 0.500% or less, Cr: 1.00% or less, Cu: 1.00% or less, and Ni: 0.50%. Hereinafter, Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.200% or less, W: 0.400% or less, Zr: 0.0200% or less, Ca: 0.0200% or less, Mg: 0.0200% or less, Co: 0.020% or less, It may contain at least one element selected from REM: 0.0200% or less, Te: 0.020% or less, Hf: 0.10% or less, and Bi: 0.200% or less.

V: 0.100 or less

V has the effect of increasing strength by forming fine carbides. If the V content exceeds 0.100%, coarse V carbides precipitate and the toughness decreases. The lower limit of the V content is not particularly limited and may be 0.000%, but it is preferably 0.001% or more because it has the effect of increasing strength by forming fine carbides.

Mo: 0.500% or less

Mo has the effect of improving quenching properties and increasing the bainite and martensite fractions. When the Mo content exceeds 0.500%, the effect is saturated. The lower limit of the Mo content is not particularly limited and may be 0.000%. However, since it has the effect of improving quenchability and increasing the bainite and martensite fractions, it is preferably set at 0.010% or more.

Cr: 1.00% or less

Cr has the effect of improving quenching properties and increasing the bainite and martensite fractions. When the Cr content exceeds 1.00%, the effect is saturated. The lower limit of the Cr content is not particularly limited and may be 0.000%. However, since it has the effect of improving quenchability and increasing the bainite and martensite fractions, it is preferable to set it to 0.01% or more.

Cu: 1.00% or less

Cu has the effect of increasing strength through solid solution. If the Cu content exceeds 1.00%, grain boundary cracks are likely to occur. The lower limit of the Cu content is not particularly limited and may be 0.000%, but since it has the effect of increasing strength through solid solution, it is preferable to set it to 0.01% or more.

Ni: 0.50% or less

Ni has the effect of improving hardenability, but the effect is saturated when the Ni content exceeds 0.50%. The lower limit of the Ni content is not particularly limited and may be 0.000%, but it is preferably 0.01% or more because it has the effect of improving quenching properties.

Sb: 0.200% or less

Sb has the effect of suppressing surface oxidation, nitriding, and decarburization of steel sheets, but the effect is saturated when the Sb content exceeds 0.200%. The lower limit of the Sb content is not particularly limited and may be 0.000%, but it is preferably 0.001% or more because it has the effect of suppressing surface oxidation, nitriding, and decarburization of the steel sheet.

Sn: 0.200% or less

Sn, like Sb, has the effect of suppressing surface oxidation, nitriding, and decarburization of steel sheets. When the Sn content exceeds 0.200%, the effect is saturated. The lower limit of the Sn content is not particularly limited and may be 0.000%, but it is preferably 0.001% or more because it has the effect of suppressing surface oxidation, nitriding, and decarburization of the steel sheet.

Ta: 0.200% or less

Ta has the effect of increasing strength by forming fine carbides. If the Ta content exceeds 0.200%, coarse Ta carbides precipitate and the toughness decreases. The lower limit of the Ta content is not particularly limited and may be 0.000%, but since it has the effect of increasing strength by forming fine carbides, it is preferable to set it to 0.001% or more.

W: 0.400% or less

W has the effect of increasing strength by forming fine carbides. If the W content exceeds 0.400%, coarse W carbides precipitate and toughness decreases. The lower limit of the W content is not particularly limited and may be 0.000%, but it is preferably 0.001% or more because it has the effect of increasing strength by forming fine carbides.

Zr: 0.0200% or less

Zr has the effect of suppressing stress concentration and improving toughness by sphericalizing the shape of the inclusions. If the Zr content exceeds 0.0200%, inclusions are formed in large quantities and toughness decreases. The lower limit of the Zr content is not particularly limited and may be 0.000%, but it is preferably set at 0.0001% or more because it has the effect of suppressing stress concentration and improving toughness by making the shape of the inclusions spherical.

Ca: 0.0200% or less

Ca can be used as a deoxidizer. If the Ca content exceeds 0.0200%, a large amount of Ca-based inclusions are generated and the toughness decreases. The lower limit of the Ca content is not particularly limited and may be 0.000%, but it is preferably 0.0001% or more because it can be used as a deoxidizing agent.

Mg: 0.0200% or less

Mg can be used as a deoxidizer. If the Mg content exceeds 0.0200%, a large amount of Mg-based inclusions are generated and the toughness decreases. The lower limit of the Mg content is not particularly limited and may be 0.000%, but it is preferably 0.0001% or more because it can be used as a deoxidizing agent.

Co: 0.020% or less

Co has the effect of increasing strength through solid solution strengthening. When the Co content exceeds 0.020%, the effect is saturated. The lower limit of the Co content is not particularly limited and may be 0.000%. However, since it has the effect of increasing strength through solid solution strengthening, it is preferable to set it to 0.001% or more.

REM: 0.0200% or less

REM has the effect of suppressing stress concentration and improving toughness by sphericalizing the shape of the inclusions. If the REM content exceeds 0.0200%, inclusions are formed in large quantities and toughness decreases. The lower limit of the REM content is not particularly limited and may be 0.000%, but it is preferably 0.0001% or more because it has the effect of suppressing stress concentration and improving toughness by making the shape of the inclusions spherical.

Te: 0.020% or less

Te has the effect of suppressing stress concentration and improving toughness by sphericalizing the shape of the inclusions. If the Te content exceeds 0.020%, inclusions are formed in large quantities and toughness decreases. The lower limit of the Te content is not particularly limited and may be 0.000%, but it is preferably 0.001% or more because it has the effect of suppressing stress concentration and improving toughness by making the shape of the inclusions spherical.

Hf: 0.10% or less

Hf has the effect of suppressing stress concentration and improving toughness by sphericalizing the shape of the inclusions. If the Hf content exceeds 0.10%, inclusions are formed in large quantities and toughness decreases. The lower limit of the Hf content is not particularly limited and may be 0.000%, but it is preferably 0.01% or more because it has the effect of suppressing stress concentration and improving toughness by making the shape of the inclusions spherical.

Bi: 0.200% or less

Bi has the effect of reducing segregation and improving bendability. If the Bi content exceeds 0.200%, inclusions are formed in large quantities and bendability deteriorates. The lower limit of the Bi content is not particularly limited and may be 0.000%, but it is preferably 0.001% or more because it has the effect of reducing segregation and improving bendability.

The remainder other than the above-mentioned components is Fe and inevitable impurities. Additionally, if the content of the above-mentioned optional elements is less than the lower limit, the effect of the present invention is not impaired. Therefore, if these optional elements are contained less than the lower limit, they are treated as unavoidable impurities.

[strong organization]

Next, the steel structure of the high-strength steel plate will be explained.

Martensite and Bainite: Total area ratio is 95% or more

Both martensite and bainite are hard phases, and are required to achieve a TS of 1180 MPa or more. Therefore, the total area ratio of martensite and bainite is set to 95% or more. The total area ratio of martensite and bainite is preferably 96% or more. The upper limit of the total area ratio of martensite and bainite is not particularly limited, and may be 100%.

The steel structure may contain residual structures other than martensite and bainite. The remaining structure includes ferrite, retained austenite, and cementite. The remaining structure is set to 5% or less in total area ratio.

Here, the area ratio of each tissue is measured as follows. The area ratio of retained austenite was determined by chemically polishing the rolled surface to a position of 1/4t of the sheet thickness of the steel sheet in the test piece collected from each steel sheet, and measuring the X of the polished surface using an X-ray diffraction (XRD) device. The line diffraction intensity and diffraction peak position are measured to calculate the volume ratio, and the number is taken as the area ratio of retained austenite. Next, the cross section of the sheet thickness parallel to the rolling direction of each steel sheet is polished and then corroded with 3% nital, and the position of 1/4t of the sheet thickness is used as the observation surface. For the observation surface, SEM images of 3 fields of view are taken at a magnification of 2000 times. For the obtained SEM image, the area ratio of the sum of martensite, bainite, and retained austenite and the area ratio of structures other than martensite, bainite, and retained austenite (ferrite, cementite) are determined through image analysis. The area ratios of martensite and bainite are obtained by subtracting the area ratio of retained austenite obtained by XRD from the area ratios of martensite, bainite, and retained austenite obtained by image analysis. The average value of the three fields of view is taken as the area ratio of the tissue.

Average grain size of old austenite grains: 10㎛ or less

By refining the crystal grain size and complicating the crack propagation path, toughness and bendability can be improved. Additionally, it has the effect of increasing yield strength by refining and strengthening the crystal grains. In order to obtain these effects, the average grain size of the prior austenite grains needs to be 10 μm or less. The average grain size of the old austenite grains is preferably 9 μm or less. The lower limit of the average grain size of the old austenite grains is not particularly limited, but is preferably 1 μm or more from the viewpoint of production technology.

Here, the average grain size of the old austenite grains is measured as follows. The cross-section of the sheet thickness parallel to the rolling direction of each steel sheet is polished and then picral etched, and the microstructure at 1/4t of the sheet thickness is photographed in 3 views at a magnification of 2000x. From the obtained tissue image, the grain size of each old austenite grain is determined by image analysis, and the average value of the three views is taken as the average grain size of the old austenite grain.

B concentration at old austenite grain boundaries: 0.10% or more in mass%

B can strengthen the grain boundaries by segregating at the old austenite grain boundaries, thereby improving toughness and bendability. If the B concentration at the old austenite grain boundary is 0.10% or more in mass%, the effect is obtained. The B concentration at the old austenite grain boundary is preferably 0.15% or more in mass%, more preferably 0.20% or more. There is no upper limit to the B concentration at the old austenite grain boundaries, but it is preferably less than 20% because it appropriately prevents hard carbon boride from precipitating on the grain boundaries and further improves toughness.

Here, the B concentration at the old austenite grain boundary is measured as follows. A needle sample is produced from the area containing the old austenite grain boundaries using the SEM-FIB (Focused Ion Beam) method. The obtained needle sample is subjected to 3DAP analysis using a 3Dimensional Atom Probe (3DAP) device (LEAP4000XSi, manufactured by AMETEK). Measurements are performed in laser mode. The B concentration at the old austenite grain boundaries is determined from the number of B ions detected from the old austenite grain boundaries and the number of other ions.

C thickening area

By strengthening the martensite grain boundaries and the parent phase sandwiching the martensite grain boundaries by C enrichment, bendability and yield ratio can be improved. In addition, in this specification, “martensite grain boundaries” includes all old austenite grain boundaries, block grain boundaries, and packet grain boundaries that exist in martensite and bainite. Figure 1 shows an example of a C-enriched region. Fig. 1(A) is a diagram showing the observation results of C-enriched regions present at block grain boundaries and packet grain boundaries. FIG. 1(B) is a diagram showing the observation results of the C-enriched region existing at the old austenite grain boundary. In Figures 1 (A) and (B), the drawings shown on the left are an example of the observation results from a scanning transmission electron microscope (STEM), and a martensite grain boundary exists in the center of the drawing. You can see that The figure shown on the right is an example of the observation results of C enrichment in STEM. From these figures, it can be seen that a C-enriched region exists along the martensite grain boundary and across the base material sandwiching the martensite grain boundary.

C concentration in C-enriched area: 4.0 times or more than the C content in steel

In the C-enriched region, sufficient grain boundary strength can be obtained by concentrating C to 4.0 times or more than the C content in the steel. That is, the C concentration in the C-enriched region satisfies the following equation (2).

C concentration (mass %) in C enriched area/C content in steel (mass %) ≥ 4.0... (2)

The C concentration in the C-enriched region is preferably 4.5 times or more than the C content in the steel. There is no particular upper limit to the C concentration in the C enriched region, but in order to appropriately prevent precipitation of cementite and appropriately prevent a decrease in the solid solution C concentration, the C concentration is preferably 6% or less.

C enrichment area: Enrichment width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary.

As shown in FIG. 1, the bendability and yield ratio can be improved by strengthening not only the martensite grain boundaries but also the parent phase sandwiching the martensite grain boundaries by C enrichment. Therefore, the C-enriched region is made to exist with an enrichment width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary. When the enrichment width of the C enrichment area is less than 3 nm, the above effect is small. On the other hand, if the width of the C-enriched region exceeds 100 nm, C cannot be sufficiently concentrated at grain boundaries and near grain boundaries. The width of the C-enriched region is preferably 3.5 nm or more. Additionally, the width of the C-enriched region is preferably 80 nm or less.

C enrichment region: length of 100 nm or more in the direction parallel to the martensite grain boundary

To obtain excellent bendability and yield ratio, it is important to strengthen the martensite grain boundaries into a network by C segregation. Therefore, the C-enriched region is made to have a length of 100 nm or more in the direction parallel to the martensite grain boundary. When the C-enriched area is less than 100 nm, destruction or yielding occurs from the nodes in the C-enriched area. The C-enriched region preferably has a length of 120 nm or more in the direction parallel to the martensite grain boundaries. There is no upper limit to the length of the C-enriched region along the martensite grain boundary, and the C-enriched region may exist so as to cover the entire length of the martensite grain boundary.

Here, the C concentration, enrichment width, and length of the C-enriched region are measured as follows. A thin film sample is created from the area containing the martensite grain boundary by the SEM-FIB method, and C plane analysis is performed by STEM and Energy Dispersive X-ray Spectroscopy (EDS). For analysis, an analytical transmission electron microscope Talos F200X (manufactured by FEI) is used. The thin film sample is tilted so that the martensite grain boundaries are parallel to the electron beam, and surface analysis is performed in an area of 200 × 500 nm. The analysis length in the direction parallel to the martensite grain boundary (direction along the martensite grain boundary) is set to 500 nm. The surface analysis data is integrated in a direction parallel to the martensite grain boundaries, and a line profile with a length of 200 nm is obtained in the direction perpendicular to the martensite grain boundaries. In the C concentration line profile, the half value of the maximum value of the line profile is obtained, and the width of the line profile that is more than half value is taken as the enrichment width of the C enrichment region. For the enrichment range, quantitative analysis of EDS is performed to quantify the C concentration in the C enrichment area. Additionally, through C plane analysis, the length of the C-enriched region is measured in a direction parallel to the martensite grain boundary, and is set as the length of the C-enriched region along the martensite grain boundary.

According to the present invention, a high-strength steel sheet with a tensile strength of 1180 MPa or more can be provided. The tensile strength of the high-strength steel plate is preferably 1250 MPa or more.

The above-mentioned high-strength steel sheet may have a plating layer on at least one side. As the plating layer, any one of a hot-dip galvanized layer, an alloyed hot-dip galvanized layer, and an electric galvanized layer is preferable. The composition of the plating layer is not particularly limited and can be a known composition.

The composition of the hot-dip galvanized layer is not particularly limited, and any general composition may be sufficient. In one example, the plating layer contains Fe: 20 mass% or less, Al: 0.001 mass% or more and 1.0 mass% or less, and further contains Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, A composition containing a total of 0% by mass or more and 3.5% by mass or less of one or two or more types selected from the group consisting of Ca, Cu, Li, Ti, Be, Bi, and REM, with the balance consisting of Zn and inevitable impurities. have When the plating layer is a hot-dip galvanized layer, in one example, the Fe content in the plating layer is less than 7% by mass, and when the plating layer is an alloyed hot-dip galvanized layer, in one example, the Fe content in the plating layer is 7% by mass or more and 15% by mass or less, More preferably, it is 8 mass% or more and 13 mass% or less.

The amount of plating is not particularly limited, but it is preferable that the amount of plating per side of the high-strength steel sheet is 20 to 80 g/m2. In one example, the plating layer is formed on both the front and back sides of a high-strength steel sheet.

Next, a method for manufacturing high-strength steel sheets will be described.

First, a steel slab having the above-described component composition is manufactured. First, the steel material is smelted to obtain molten steel having the above composition. The solvent method is not particularly limited, and any known solvent method, such as a converter solvent or an electric furnace solvent, is suitable. The obtained molten steel is solidified to produce a steel slab. The method of manufacturing a steel slab from molten steel is not particularly limited, and continuous casting, ingot method, or thin slab casting method can be used. The steel slab may be cooled once and then heated again before hot rolling, or the cast steel slab may be continuously hot rolled without cooling to room temperature. The slab heating temperature is preferably set to 1100°C or higher, and is preferably set to 1300°C or lower, considering rolling load and scale generation. The method of heating the slab is not particularly limited, but for example, it can be heated in a heating furnace according to a conventional method.

[Hot rolling process]

Next, the heated steel slab is hot rolled to form a hot rolled sheet. There are no particular restrictions on hot rolling, and it may be performed according to commercial methods. Cooling after hot rolling is also not particularly limited, and it is cooled to the coiling temperature. Next, the hot rolled sheet is wound into a coil. The coiling temperature is preferably 400°C or higher. This is because if the coiling temperature is 400°C or higher, coiling becomes easy without increasing the strength of the hot-rolled sheet. The coiling temperature is more preferably 550°C or higher. In addition, in order to appropriately prevent the formation of thick scale and further improve the yield, the coiling temperature is preferably set to 750°C or lower. Additionally, before acid cleaning, heat treatment may be performed on the hot-rolled sheet for the purpose of softening it.

[Acid washing process]

Optionally, the scale of the hot rolled sheet wound into a coil is removed. The method for removing scale is not particularly limited, but in order to completely remove scale, it is preferable to perform acid cleaning while rewinding the hot rolled coil. The acid washing method is not particularly limited, and commercial methods may be used.

[Cold rolling process]

After the hot-rolled sheet from which scale has been arbitrarily removed is properly cleaned, it is cold-rolled to obtain a cold-rolled sheet. The method of cold rolling is not particularly limited, and commercial methods may be used.

[Annealing process]

Next, the cold-rolled sheet is heated to a first heating temperature of 850°C or more and 920°C or less and maintained for 10 s or more, and then the temperature is raised to a second heating temperature of 1000°C or more and 1200°C or less at an average heating rate of 50°C/s or more. Then, an annealing process is performed within 5 seconds after reaching the second heating temperature, cooling to 500°C or lower at a cooling rate of 50°C/s or more.

First heating temperature of 850℃ or higher and 920℃ or lower

Next, the cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and maintained for 10 seconds or longer. In order to obtain martensite and bainite-based structures, annealing is performed at the first heating temperature in the austenite single phase region. If the first heating temperature is less than 850°C, ferrite is formed and the strength decreases. On the other hand, when the first heating temperature exceeds 920°C, the austenite grain size exceeds 10 μm and cannot be refined in subsequent processes, thereby reducing bendability, toughness, and yield ratio. The first heating temperature is preferably 860°C or higher. Moreover, the first heating temperature is preferably 900°C or lower.

Preservation time at first heating temperature: 10 s or more

The preservation time at the first heating temperature is 10 s or more. By holding the material at the first heating temperature for 10 seconds or more, the growth of austenite grain size and the growth inhibition due to pinning or solid solution by Nb carbide are balanced. If the holding time is less than 10 s, the austenite grains are in the middle of growth, and the effect of pinning by Nb carbide or inhibiting growth by solid solution does not occur during continued rapid heating, and the old austenite grain size exceeds 10 μm. The upper limit of the preservation time at the first heating temperature is not particularly limited, but from the viewpoint of productivity, the preservation time at the first heating temperature is preferably 60 s or less. The preservation time at the first heating temperature is preferably 20 s or more.

Second heating temperature of 1000℃ or higher and 1200℃ or lower

After preservation at the first heating temperature, annealing is performed at a high temperature while maintaining the austenite grain boundaries at 10 μm or less to segregate a sufficient amount of B at the grain boundaries. If the second heating temperature is less than 1000°C, diffusion of B is slow and grain boundary segregation is insufficient. When the second heating temperature exceeds 1200°C, the growth of austenite grains is rapid, and the austenite grain size exceeds 10 μm. The second heating temperature is preferably 1020°C or higher. The second heating temperature is preferably 1150°C or lower.

Average heating rate: more than 50℃/s

The average heating rate from the first heating temperature to the second heating temperature is 50°C/s or more. When the average heating rate from the first heating temperature to the second heating temperature is less than 50°C/s, the austenite grain size grows to more than 10 μm. The upper limit of the average heating rate from the first heating temperature to the second heating temperature is not particularly limited, but is preferably 120°C/s or less because excessively rapid heating is difficult to control. The average heating rate from the first heating temperature to the second heating temperature is preferably 80°C/s or more.

Cooling at an average cooling rate of 50℃/s or more to 500℃ or less within 5 seconds after reaching the second heating temperature

After reaching the second heating temperature, without maintaining it at the second heating temperature, rapid cooling is started within 5 seconds after reaching the second heating temperature, and rapid cooling is performed to 500°C or lower at an average cooling rate of 50°C/s or more. Do. Accordingly, it is possible to obtain a steel structure in which 0.1% or more of B is grain boundary segregated at an austenite grain size of 10 μm or less. When maintained at the second heating temperature, grain growth begins quickly, so cooling is started immediately after reaching the second heating temperature.

Average cooling rate: more than 50℃/s

In cooling after reaching the second heating temperature, the average cooling rate from the second heating temperature to 500°C or lower is 50°C/s or more. If the average cooling rate from the second heating temperature to 500°C or lower is less than 50°C/s, grain growth occurs during cooling. The upper limit of the average cooling rate from the second heating temperature to 500°C or lower is not particularly limited, but is preferably 120°C/s or lower to facilitate control. The average cooling rate from the second heating temperature to 500°C or lower is preferably 80°C/s or higher.

Cooling stop temperature: below 500℃

Additionally, in order to suppress ferrite transformation, rapid cooling is performed to a cooling stop temperature of 500°C or lower. The cooling stop temperature is preferably 450°C or lower. The lower limit of the cooling stop temperature is not particularly limited, but is preferably 100°C or higher.

After the above-described annealing process and before the reheating process, a plating process may be performed by performing a plating treatment on at least one side of the high-strength steel sheet to obtain a high-strength plated steel sheet. After the plating process, the high-strength plated steel sheet may be subjected to heat treatment to alloy the plating layer of the high-strength plated steel sheet, and an alloyed plated steel sheet may be obtained.

After the annealing process, rolling process is carried out with an elongation of 0.5% or more.

After the annealing process described above, a rolling process is performed in which the cold-rolled sheet is rolled to an elongation of 0.5% or more to obtain a second cold-rolled sheet. The cold-rolled sheet obtained through the processes up to this point contains many moving dislocations. Through this rolling process, movable dislocations accumulate at the grain boundaries and become entangled with each other to become immobile dislocations. If the elongation rate is less than 0.5%, the effect is small. The elongation rate in the rolling process is preferably 0.6% or more. There is no particular upper limit to the elongation rate in the rolling process, but in order to further reduce the load on equipment, for example, 2% or less is preferable.

After the rolling process, a reheating process of maintaining the second cold rolled sheet at a reheating temperature of 70°C or more and 200°C or less for 600s or more.

After the above-described rolling process, the second cold-rolled sheet is tempered at a low temperature in order to segregate C on dislocations accumulated near grain boundaries or generate clusters. If the reheating temperature is less than 70°C, diffusion of C is slow, and C does not concentrate near the grain boundaries to a sufficient amount. On the other hand, if the reheating temperature exceeds 200°C, excessive tempering proceeds and cementite precipitates. Cementite precipitated at the grain boundary easily becomes a fracture origin, and the C concentration of the matrix around the cementite decreases, resulting in a decrease in bendability and toughness. The reheating temperature is preferably 90°C or higher. Additionally, the reheating temperature is preferably 190°C or lower.

Preservation time at reheating temperature: 600s or more

If the storage time at the reheating temperature is less than 600 s, diffusion of C is slow, and a sufficient amount of C concentration cannot be obtained. The upper limit of the storage time at the reheating temperature is not particularly limited, but is preferably 43200 s (0.5 day) or less to prevent precipitation of cementite. The preservation time at the reheating temperature is preferably 800 s or more.

In addition, when reheating is performed without performing the above-mentioned rolling process, C segregates at the grain boundaries and toughness improves, but because the concentration range is narrow and no strengthening occurs except at the grain boundaries, bendability is inferior. Additionally, the potential remains at the movable potential and YR lags behind.

Additionally, manufacturing conditions other than those described above may follow commercial methods.

[absence]

It is possible to provide a member made of at least part of the above-described high-strength steel sheet or high-strength plated steel sheet. In one example, the above-described high-strength steel sheet or high-strength plated steel sheet can be formed into a desired shape by press working to make automobile parts. Additionally, the automobile parts may contain steel sheets other than the high-strength steel sheet or high-strength plated steel sheet according to the present embodiment as a material. According to this embodiment, it is possible to provide a high-strength steel sheet with a TS of 1180 MPa or more and which has bendability, toughness, and a high yield ratio. Therefore, it is suitable as an automobile part that contributes to reducing the weight of the automobile body. This high-strength steel sheet or high-strength plated steel sheet can be suitably used in all automotive parts, especially members used as skeletal structural parts or reinforcement parts.

Example

Steel with the component composition shown in Table 1 and the balance consisting of Fe and inevitable impurities was melted in a converter to make a steel slab. The obtained slab was reheated, hot rolled, and wound to obtain a hot rolled coil. Next, acid cleaning treatment was performed while rewinding the hot rolled coil, and cold rolling was performed. The sheet thickness of the hot-rolled sheet was 3.0 mm, and the sheet thickness of the cold-rolled sheet was 1.2 mm. Annealing was performed under the conditions shown in Table 2 in a continuous hot-dip galvanizing line to obtain cold-rolled steel sheets, hot-dip galvanized steel sheets (GI), and alloyed hot-dip galvanized steel sheets (GA). The hot-dip galvanized steel sheet was immersed in a plating bath at 460°C, and the plating adhesion amount per side was 35 g/m2. The alloyed hot-dip galvanized steel sheet was manufactured by adjusting the plating amount to 45 g/m2 per side and then performing alloying treatment with storage at 520°C for 40 s. The obtained steel sheet was subjected to rolling processing and reheating under the conditions shown in Table 2.

For the obtained steel sheet, according to the method described above, the sum of the area ratios of martensite and bainite, the old austenite grain size, the B concentration at the old austenite grain boundaries, and the C concentration (mass%) in the C-enriched region at the martensite grain boundaries. )/C content (mass%) in the steel, enrichment width of the C-enriched region, and length along the martensite grain boundary of the C-enriched region were evaluated. In addition, tensile strength, yield ratio, toughness, and bendability were evaluated according to the method described later. The results are shown in Table 3.

[Tensile test]

A tensile test was performed on the obtained steel sheet in accordance with JIS Z 2241. A JIS5 tensile test piece was collected with the direction perpendicular to the rolling direction as the longitudinal direction, and a tensile test was performed to measure the tensile strength (TS) and yield strength (YS). If the tensile strength TS was 1180 MPa or more, the tensile strength was judged to be good. Additionally, if the ratio YR=YS/TS of yield strength and tensile strength was 0.80 or more, it was considered a high yield ratio.

[Charpy test]

The Charpy impact test was conducted in accordance with JIS Z 2242. From the obtained steel sheet, a 90° V notch was created so that the direction perpendicular to the rolling direction of the steel sheet was the V notch application direction, the width was 10 mm, the length was 55 mm, and the notch depth was 2 mm in the center of the length. collected. After that, a Charpy impact test was performed in a test temperature range of -120 to +120°C. A transition curve was obtained from the obtained brittle fracture ratio, and the temperature at which the brittle fracture ratio was 50% was determined to be the brittle-ductile transition temperature. In addition, toughness was judged to be good when the brittle-ductile transition temperature obtained from the Charpy test was -40°C or lower. In the table, when the brittle-ductile transition temperature is -40°C or lower, the toughness is indicated as "excellent," and when the brittle-ductile transition temperature is above -40°C, the toughness is indicated as "poor."

[Bending test]

The bending test was performed in accordance with JIS Z 2248. From the obtained steel sheet, a rectangular test piece with a width of 30 mm and a length of 100 mm was taken so that the direction parallel to the rolling direction of the steel sheet was the axial direction of the bending test. After that, a 90°V bending test was performed under the conditions of a press load of 100 kN and a press hold time of 5 seconds. In addition, bendability is evaluated by the passing rate of the bending test, and the maximum R at which the bending radius (R) divided by the plate thickness (t), R/t, is 5 or less (for example, when the plate thickness is 1.2 mm, At a bending radius of 7.0 mm), a bending test is performed on 5 samples, and then the presence or absence of cracks in the ridge portion of the bending apex is evaluated. If all 5 samples are not cracked, that is, the pass rate is 100%. However, it was judged that the bendability was good. In the table, only cases where the passing rate was 100% were indicated as “excellent” in bendability, and other cases were indicated as “poor” in bendability. Here, the presence or absence of cracks was evaluated by measuring the ridge portion of the bending peak at a magnification of 40 times using a digital microscope (RH-2000: manufactured by Hyrox Co., Ltd.).

From Table 3, in the present invention example, TS is 1180 MPa or more, yield ratio is 0.80 or more, and bendability and toughness are excellent. On the other hand, in the comparative example, one or more of TS, yield ratio, bendability, and toughness were inferior.

Claims (7)

In mass%,
C: 0.10% or more and 0.30 or less,
Si: 0.20% or more and 1.20% or less,
Mn: 2.5% or more and 4.0% or less,
P: 0.050% or less,
S: 0.020% or less,
Al: 0.10% or less,
N: 0.01% or less,
Ti: 0.100% or less,
Nb: 0.002% or more and 0.050% or less and
B: 0.0005% or more and 0.0050% or less
Contains, the balance consists of Fe and inevitable impurities, and has a component composition that satisfies the following formula (1),
The total area ratio of martensite and bainite is 95% or more,
The average grain size of the old austenite grains is 10㎛ or less,
The B concentration of the old austenite grain boundaries is 0.10% or more in mass%,
There is a C-enriched region along the martensite grain boundary,
The C concentration in the C-enriched area is 4.0 times or more than the C content in the steel,
A high-strength steel sheet having a thickening width of 3 nm or more and 100 nm or less in a direction perpendicular to the martensite grain boundaries and a length of 100 nm or more in a direction parallel to the martensite grain boundaries.
([％N]/14)/([％Ti]/47.9)＜1.0… (One)
In formula (1), [%N] and [%Ti] represent the content (% by mass) of N and Ti in the steel, respectively. According to paragraph 1,
The above component composition is further expressed in mass%,
V: 0.100 or less,
Mo: 0.500% or less,
Cr: 1.00% or less,
Cu: 1.00% or less,
Ni: 0.50% or less,
Sb: 0.200% or less,
Sn: 0.200% or less,
Ta: 0.200% or less,
W: 0.400% or less,
Zr: 0.0200% or less,
Ca: 0.0200% or less,
Mg: 0.0200% or less,
Co: 0.020% or less,
REM: 0.0200% or less,
Te: 0.020% or less,
Hf: 0.10% or less and
Bi: 0.200% or less
A high-strength steel sheet containing at least one element selected from the group consisting of: A high-strength plated steel sheet having a plating layer on at least one side of the high-strength steel sheet according to claim 1 or 2. A steel slab having the composition of claim 1 or 2 is subjected to hot rolling to obtain a hot rolled sheet,
Cold rolling is performed on the hot-rolled sheet to obtain a cold-rolled sheet,
The cold-rolled sheet is heated to a first heating temperature of 850°C or more and 920°C or less and held for 10 s or more, and then heated to a second heating temperature of 1000°C or more and 1200°C or less at an average heating rate of 50°C/s or more. An annealing process is performed in which the temperature is raised to and cooled to 500°C or lower at an average cooling rate of 50°C/s or more within 5 seconds after reaching the second heating temperature,
After the annealing process, a rolling process is performed to obtain a second cold-rolled sheet by rolling the cold-rolled sheet to an elongation of 0.5% or more,
A method of manufacturing a high-strength steel sheet, in which, after the rolling process, a reheating process is performed in which the second cold-rolled sheet is maintained at a reheating temperature of 70°C or higher and 200°C or lower for 600s or more to obtain a high-strength steel plate. A method for producing a high-strength plated steel sheet, comprising a plating process of obtaining a high-strength plated steel sheet by performing a plating treatment on at least one side of the high-strength steel plate after the annealing process according to claim 4 and before the reheating process. A member formed by using at least part of the high-strength steel plate according to claim 1 or 2. A member formed by using at least part of the high-strength plated steel sheet according to claim 3.