High-strength steel sheet, high-strength plated steel sheet, and method for manufacturing same, and component

Abstract

The present invention provides a high-strength steel sheet having a tensile strength of 1180 MPa or more and excellent bendability and toughness, and a manufacturing method thereof. The high-strength steel sheet of the present invention has a composition containing C, Si, Mn, P, S, Al, N, Ti, Nb, B, and the remainder consisting of Fe and inevitable impurities, satisfies the following formula (1), and has a total area ratio of martensite and bainite of 95% or more, an average crystal grain size of prior austenite grains of 10 μm or less, a B concentration at a prior austenite grain boundary of 0.10% or more in mass%, a C-rich region along a martensite grain boundary, the C-rich region having a C concentration of 4.0 times or more the C content in the steel, a concentration width of 3 nm to 100 nm in a direction orthogonal to the martensite grain boundary, and a length of 100 nm or more in a direction parallel to the martensite grain boundary. ([%N] / 14) / ([%Ti] / 47.9) < 1.0 … (1).

Description

Technical Field

[0001] This invention relates to a high-strength steel plate, its manufacturing method, and components. Background Technology

[0002] For automotive steel sheets, high strength is required to improve fuel efficiency through lightweight vehicle bodies. The frame components need high-strength steel sheets with a tensile strength of 1180 MPa or higher. Furthermore, high flexibility is required to press the steel sheets into desired shapes. Moreover, from a crash safety perspective, in addition to strength, automotive components must ensure that the occupant's seating space does not easily deform during a collision. For such components, steel sheets with a high yield ratio are desirable. In addition, high toughness is also required to prevent the components from fracturing during a collision.

[0003] Patent Document 1 discloses a high-strength steel sheet 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 the same. Patent Document 3 discloses a high-yield-ratio type high-strength steel sheet and a method for manufacturing the same.

[0004] Existing technical documents

[0005] Patent documents

[0006] Patent Document 1: Japanese Patent No. 5728108

[0007] Patent Document 2: Japanese Patent No. 6597939

[0008] Patent Document 3: Japanese Patent No. 6700398 Summary of the Invention

[0009] However, yield ratio was not considered in patent documents 1 and 2. Toughness was not considered in patent document 3.

[0010] It is difficult to manufacture high-strength steel plates with a tensile strength of over 1180 MPa and excellent bending and toughness in the current technology.

[0011] The present invention was made in view of the above circumstances, and its object is to provide a high-strength steel plate with a high yield ratio, having a tensile strength of 1180 MPa or more and possessing excellent bending and toughness, and a method for manufacturing the same.

[0012] It should be noted that in this invention, high strength refers to a tensile strength TS of 1180 MPa or higher as measured according to JIS Z2241.

[0013] In addition, excellent bending performance means that no breakage occurs at the edge of the bend apex during a bending test conducted in accordance with JIS Z2248.

[0014] In addition, excellent toughness means that the brittle-ductile transition temperature is below -40°C in the Charpy impact test conducted according to JIS Z2242.

[0015] In addition, a high yield ratio refers to a yield strength to tensile strength ratio YS / TS of 0.80 or higher as determined by JIS Z 2241.

[0016] In order to achieve the above-mentioned problem, the inventors have repeatedly conducted in-depth research and obtained the following insights.

[0017] (1) The crack initiation, propagation, and brittle fracture path during bending are along the original austenite grain boundaries. Therefore, refining the grain size to complicate the fracture path and increasing the strength of the grain boundaries are effective in improving bending performance. For refining the original austenite grains, it is effective to minimize the annealing temperature above 850°C, which is the austenite single-phase region. On the other hand, for strengthening the grain boundaries, it is effective to induce boron grain boundary segregation, but the amount of boron grain boundary segregation increases with annealing at higher temperatures. Therefore, in order to increase the amount of boron grain boundary segregation while maintaining a fine grain size, rapid heating and rapid cooling are performed after annealing at around 850°C to obtain fine austenite grains. Thus, while suppressing grain growth, it is possible to promote boron grain boundary segregation caused by diffusion, thereby simultaneously achieving austenite grain size refinement and boron grain boundary segregation.

[0018] (2) The dislocations present in the martensite structure after quenching are movable dislocations that are prone to slip movement under low stress, thus the yield stress of the martensite structure is low. However, if the quenched steel plate is slightly processed, these dislocations move towards the grain boundaries and become immobile dislocations. This can increase the yield ratio of the steel plate.

[0019] (3) By tempering the steel plate at a low temperature, carbon segregates or precipitates on the dislocations. By tempering the processed steel plate, where dislocations have accumulated near the grain boundaries, a region with high C concentration (C-enriched region) is formed along the grain boundary network, resulting in a significant increase in strength near the grain boundaries. C is enriched not only at the grain boundaries but also in the parent phase sandwiched between the grain boundaries, thus greatly increasing its strength. Since grain boundaries are not easily deformed, the yield ratio also increases significantly by forming C-enriched regions.

[0020] This invention is based on the above-mentioned insights. Specifically, the main elements of this invention are as follows.

[0021] [1] A high-strength steel plate has the following composition: by mass % containing C: 0.10% to 0.30%, Si: 0.20% to 1.20%, Mn: 2.5% to 4.0%, P: less than 0.050%, S: less than 0.020%, Al: less than 0.10%, N: less than 0.01%, Ti: less than 0.100%, Nb: 0.002% to 0.050%, and B: less than 0.0005%; the remainder consists of Fe and unavoidable impurities, and satisfies the following formula (1).

[0022] Furthermore, the combined area ratio of martensite and bainite is over 95%.

[0023] The average grain size of the original austenite grains is less than 10 μm.

[0024] The boron concentration at the original austenite grain boundaries is greater than 0.10% by mass.

[0025] Along the martensite grain boundaries, there are C-rich regions.

[0026] The C concentration in the aforementioned C-enriched region is more than 4.0 times that of the C content in the steel.

[0027] In a direction orthogonal to the aforementioned martensite grain boundary, it has an enrichment width of 3 nm to 100 nm, and in a direction parallel to the aforementioned martensite grain boundary, it has a length of more than 100 nm.

[0028] ([%N] / 14) / ([%Ti] / 47.9)<1.0…(1)

[0029] In equation (1), [%N] and [%Ti] represent the N and Ti content (mass %) in the steel, respectively.

[0030] [2] According to the high-strength steel plate described in [1] above, wherein the above composition further contains, by mass %, at least one element selected from V: less than 0.100%, Mo: less than 0.500%, Cr: less than 1.00%, Cu: less than 1.00%, Ni: less than 0.50%, Sb: less than 0.200%, Sn: less than 0.200%, Ta: less than 0.200%, W: less than 0.400%, Zr: less than 0.0200%, Ca: less than 0.0200%, Mg: less than 0.0200%, Co: less than 0.020%, REM: less than 0.0200%, Te: less than 0.020%, Hf: less than 0.10%, and Bi: less than 0.200%.

[0031] [3] A high-strength galvanized steel sheet having a galvanized layer on at least one side of the high-strength steel sheet described in [1] or [2] above.

[0032] [4] A method for manufacturing a high-strength steel plate, comprising the following steps:

[0033] Hot-rolled plates are produced by hot rolling steel slabs having the composition described in [1] or [2] above.

[0034] The above-mentioned hot-rolled sheet is cold-rolled to produce cold-rolled sheet.

[0035] The cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and held for at least 10 seconds. Then, it is heated to a second heating temperature of 1000°C or higher and 1200°C or lower at an average heating rate of 50°C / s or higher. After reaching the second heating temperature, it is cooled to below 500°C at an average cooling rate of 50°C / s or higher within 5 seconds, thereby performing an annealing process.

[0036] Following the annealing process described above, a rolling process is performed to roll the cold-rolled sheet with an elongation of 0.5% or more to obtain a second cold-rolled sheet.

[0037] After the above rolling process, a reheating process is performed to hold the second cold-rolled sheet at a reheating temperature of 70°C or higher and 200°C or lower for 600 seconds or more to obtain a high-strength steel sheet.

[0038] [5] A method for manufacturing a high-strength coated steel sheet, wherein, after the annealing process described in [4] and before the reheating process, a coating process is performed on at least one side of the high-strength steel sheet to obtain a high-strength coated steel sheet.

[0039] [6] A component, at least a portion of which is made of the high-strength steel plate described in [1] or [2] above.

[0040] [7] A component, at least a portion of which is made of the high-strength plated steel sheet described above [3].

[0041] Invention Effects

[0042] According to the present invention, a high-strength steel plate with a high yield ratio, possessing a tensile strength of 1180 MPa or more and excellent bending and toughness, and a method thereof are provided. Attached Figure Description

[0043] Figure 1 This is a diagram representing an example of a C-rich region. Detailed Implementation

[0044] The following describes embodiments of the present invention. It should be noted that the present invention is not limited to the following embodiments. First, a suitable range for the composition of the steel plate and the reasons for this limitation will be explained. It should be noted that in the following description, the "%" indicating the content of the constituent elements in the steel plate refers to "mass %" unless otherwise specified. Furthermore, in this specification, the numerical range indicated by "~" refers to the range of values ​​before and after "~" as a lower and upper limit.

[0045] C: 0.10%~0.30

[0046] C has the following effects: besides strengthening martensite and bainite structures, it also segregates to dislocations that accumulate near the original austenite grain boundaries, thus strengthening the grain boundaries and improving flexibility, toughness, and yield ratio. When the C content is below 0.10%, the area fraction of martensite and bainite decreases, and a yield strength (TS) of 1180 MPa or higher cannot be obtained. When the C content exceeds 0.30%, boron carbides with iron are formed during annealing, making it impossible for a sufficient amount of boron to segregate at the original austenite grain boundaries. The C content is preferably set to 0.11% or more. Furthermore, the C content is preferably set to 0.28% or less.

[0047] Si: 0.20%~1.20%

[0048] Si is an effective element for solid solution strengthening and needs to be added at least 0.20%. On the other hand, Si is an element that stabilizes ferrite and raises the phase transformation point. Therefore, when the Si content exceeds 1.20%, it is difficult to achieve a pre-austenite grain size of less than 10 μm. The Si content is preferably set at 0.50% or more. The Si content is preferably set at 1.10% or less.

[0049] Mn: 2.5%–4.0%

[0050] Mn is effective in improving hardenability. When the Mn content is below 2.5%, the area ratio of martensite and bainite decreases, resulting in reduced strength. On the other hand, when the Mn content exceeds 4.0%, the segregated areas become excessively hardened, reducing flexibility. The Mn content is preferably set to 2.8% or more. The Mn content is preferably set to 3.5% or less.

[0051] P: below 0.050%

[0052] Phosphorus (P) segregates at the original austenite grain boundaries, reducing toughness; therefore, the P content is set to 0.050% or less. There is no specific lower limit for the P content; it can be 0%, but setting it below 0.001% would increase manufacturing costs, so it is preferably 0.001% or more. The P content is preferably set to 0.025% or less.

[0053] S: below 0.020%

[0054] Sulfur segregates at the original austenite grain boundaries, reducing toughness; therefore, the sulfur content is set to 0.020% or less. While no specific lower limit is set for the sulfur content, setting it below 0.0001% would increase manufacturing costs; therefore, it is preferable to set it to 0.0001% or more. The sulfur content is preferably set to 0.018% or less.

[0055] Al: below 0.10%

[0056] Al is the element that functions as a deoxidizer, and to achieve this effect, the Al content is preferably set at 0.005% or more. On the other hand, when the Al content exceeds 0.10%, ferrite is easily formed, resulting in reduced strength. Therefore, the Al content is preferably set at 0.05% or less.

[0057] N: less than 0.01%

[0058] N forms nitrides with Nb and B, reducing the effectiveness of Nb and B addition. Therefore, it is set to 0.01% or less. No specific lower limit is set, but from a manufacturing cost perspective, it is preferable to set it to 0.0001% or more.

[0059] Ti: below 0.100%

[0060] Ti fixes nitrogen in steel in the form of TiN, which has the effect of suppressing the formation of BN and NbN. To obtain these effects, the Ti content is preferably set to 0.005% or more. On the other hand, when the Ti content exceeds 0.100%, coarse Ti carbides form at the grain boundaries, reducing toughness. The Ti content is preferably set to 0.05% or less.

[0061] Nb: 0.002%~0.050%

[0062] Nb is dissolved or precipitated as fine carbides, inhibiting the growth of austenite grains during annealing. To achieve this effect, the Nb content is set to 0.002% or more. On the other hand, when the Nb content exceeds 0.050%, not only does the effect saturate, but coarse Nb carbides also precipitate, reducing toughness. The Nb content is preferably set to 0.005% or more. Furthermore, the Nb content is preferably 0.040% or less.

[0063] B: 0.0005%~0.0050%

[0064] Boron (B) has the effect of increasing grain boundary strength by segregating at the original austenite grain boundaries. To achieve this effect, the B content is set to 0.0005% or more. On the other hand, when the B content exceeds 0.0050%, carborides are formed, and toughness decreases. The B content is preferably set to 0.0010% or more. Furthermore, the B content is preferably set to 0.0030% or less.

[0065] ([%N] / 14) / ([%Ti] / 47.9)<1.0…(1)

[0066] To achieve the aforementioned effects of adding B and Nb, N, which readily combines with these elements, needs to be fixed by Ti. Therefore, the mole fraction of N is made less than the mole fraction of Ti. That is, the N and Ti content in the steel is adjusted in a manner that satisfies the above equation (1). It should be noted that in equation (1), [%N] and [%Ti] represent the N and Ti content in the steel (mass %), respectively.

[0067] [Any ingredients]

[0068] In addition to the composition described above, the high-strength cold-rolled steel sheet of this embodiment may also contain, by mass percent, at least one element selected from 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.

[0069] V: below 0.100

[0070] V has the effect of forming fine carbides, thereby increasing strength. When the V content exceeds 0.100%, coarse V carbides precipitate, reducing toughness. There is no particular limit to the lower limit of V content, which can be 0.000%, but since it has the effect of forming fine carbides and increasing strength, it is preferable to set it to 0.001% or more.

[0071] Mo: 0.500% or less

[0072] Mo has the effect of improving hardenability and increasing the proportion of bainite and martensite. The effect saturates when the Mo content exceeds 0.500%. There is no particular limit to the lower limit of the Mo content, which can be 0.000%, but it is preferable to set it to 0.010% or more to achieve the effect of improving hardenability and increasing the proportion of bainite and martensite.

[0073] Cr: less than 1.00%

[0074] Cr improves hardenability and increases the proportion of bainite and martensite. The effect saturates when the Cr content exceeds 1.00%. There is no specific lower limit for the Cr content; it can be 0.000%, but considering its effects on improving hardenability and increasing the proportion of bainite and martensite, it is preferable to set it to 0.01% or higher.

[0075] Cu: below 1.00%

[0076] Cu has the effect of increasing strength through solid solution. When the Cu content exceeds 1.00%, grain boundary cracking is likely to occur. There is no particular limit to the lower limit of Cu content, which can be 0.000%, but since it has the effect of increasing strength through solid solution, it is preferred to set it to 0.01% or more.

[0077] Ni: below 0.50%

[0078] Ni improves hardenability, but the effect saturates when the Ni content exceeds 0.50%. There is no particular limit to the lower limit of Ni content, which can be 0.000%, but for the purpose of improving hardenability, it is preferable to set it to 0.01% or higher.

[0079] Sb: below 0.200%

[0080] Sb has the effect of inhibiting surface oxidation, nitriding, and decarburization of steel plates, but the effect saturates when the Sb content exceeds 0.200%. There is no particular limit to the lower limit of Sb content, which can be 0.000%, but from the perspective of inhibiting surface oxidation, nitriding, and decarburization of steel plates, it is preferred to set it to 0.001% or higher.

[0081] Sn: below 0.200%

[0082] Like Sb, Sn has the effect of inhibiting surface oxidation, nitriding, and decarburization of steel plates. The effect saturates when the Sn content exceeds 0.200%. There is no particular limit to the lower limit of Sn content, which can be 0.000%, but from the perspective of inhibiting surface oxidation, nitriding, and decarburization of steel plates, it is preferable to set it to 0.001% or more.

[0083] Ta: below 0.200%

[0084] Ta has the effect of forming fine carbides, thereby increasing strength. When the Ta content exceeds 0.200%, coarse Ta carbides precipitate, reducing toughness. There is no particular limit to the lower limit of Ta content, which can be 0.000%, but since it has the effect of forming fine carbides and increasing strength, it is preferred to set it to 0.001% or more.

[0085] W: below 0.400%

[0086] W has the effect of forming fine carbides, thereby increasing strength. When the W content exceeds 0.400%, coarse W carbides precipitate, reducing toughness. There is no particular limit to the lower limit of the W content, which can be 0.000%, but since it has the effect of forming fine carbides and increasing strength, it is preferable to set it to 0.001% or more.

[0087] Zr: below 0.0200%

[0088] Zr has the effect of spherizing inclusions, suppressing stress concentration, and improving toughness. When the Zr content exceeds 0.0200%, inclusions form in large quantities, and toughness decreases. There is no particular limit to the lower limit of Zr content, which can be 0.000%, but from the perspective of having the effect of spherizing inclusions, suppressing stress concentration, and improving toughness, it is preferable to set it to 0.0001% or more.

[0089] Ca: below 0.0200%

[0090] Ca can be used as a deoxidizing material. When the Ca content exceeds 0.0200%, a large number of Ca-based inclusions are formed, resulting in reduced toughness. There is no particular limit to the lower limit of the Ca content, which can be 0.000%, but since it can be used as a deoxidizing material, it is preferred to set it to 0.0001% or higher.

[0091] Mg: below 0.0200%

[0092] Mg can be used as a deoxidizing material. When the Mg content exceeds 0.0200%, a large number of Mg-based inclusions are formed, resulting in reduced toughness. There is no particular limit to the lower limit of the Mg content, which can be 0.000%, but since it can be used as a deoxidizing material, it is preferred to set it to 0.0001% or higher.

[0093] Co: less than 0.020%

[0094] Co has the effect of increasing strength through solid solution strengthening. The effect saturates when the Co content exceeds 0.020%. There is no particular limit to the lower limit of the Co content, which can be 0.000%, but since it has the effect of increasing strength through solid solution strengthening, it is preferred to set it to 0.001% or more.

[0095] REM: below 0.0200%

[0096] REM has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. When the REM content exceeds 0.0200%, a large number of inclusions are formed, and the toughness decreases. There is no particular limit to the lower limit of REM content, which can be 0.000%, but from the perspective of having the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness, it is preferable to set it to 0.0001% or more.

[0097] Te: below 0.020%

[0098] Te has the effect of spherizing inclusions, thereby suppressing stress concentration and improving toughness. When the Te content exceeds 0.020%, inclusions form in large quantities, and toughness decreases. There is no particular limit to the lower limit of the Te content, which can be 0.000%, but from the perspective of having the effect of spherizing inclusions, suppressing stress concentration, and improving toughness, it is preferable to set it to 0.001% or more.

[0099] Hf: below 0.10%

[0100] Hf has the effect of spherizing inclusions, thereby suppressing stress concentration and improving toughness. When the Hf content exceeds 0.10%, inclusions form in large quantities, and toughness decreases. There is no particular limit to the lower limit of Hf content, which can be 0.000%, but considering the effect of spherizing inclusions, suppressing stress concentration, and improving toughness, it is preferable to set it to 0.01% or more.

[0101] Bi: below 0.200%

[0102] Bi has the effect of reducing segregation and improving flexibility. When the Bi content exceeds 0.200%, a large number of inclusions are formed, and flexibility decreases. There is no particular limit to the lower limit of Bi content, which can be 0.000%, but from the perspective of reducing segregation and improving flexibility, it is preferable to set it to 0.001% or more.

[0103] The remainder other than the above-mentioned components consists of Fe and unavoidable impurities. It should be noted that any of the above-mentioned components, if present in amounts below the lower limit, will not impair the effectiveness of the invention; however, if the content of any of these elements is below the lower limit, they are treated as unavoidable impurities.

[0104] [Steel Structure]

[0105] Next, the microstructure of high-strength steel plates will be explained.

[0106] Martensite and bainite: The combined area ratio is over 95%.

[0107] Both martensite and bainite are hard phases, and are necessary to achieve a strength (TS) of 1180 MPa or higher. Therefore, the combined area ratio of martensite and bainite is set to 95% or higher. Preferably, the combined area ratio of martensite and bainite is 96% or higher. There is no particular upper limit to the combined area ratio of martensite and bainite, and it can be 100%.

[0108] The microstructure of steel may include residual microstructure other than martensite and bainite. Examples of residual microstructure include ferrite, retained austenite, and cementite. The residual microstructure is defined as a total area ratio of less than 5%.

[0109] Here, the area ratios of each microstructure were determined as follows. Regarding the area ratio of retained austenite, in the test pieces collected from each steel plate, the rolled surface was chemically ground to 1 / 4 t of the plate thickness. The X-ray diffraction intensity and diffraction peak positions of the ground surface were measured using an X-ray diffraction (XRD) apparatus to calculate the volume ratio, which was then used as the area ratio of retained austenite. Next, the plate thickness section parallel to the rolling direction of each steel plate was ground and etched with a 3% nitric acid ethanol solution, with the 1 / 4 t thickness position used as the observation surface. For the observation surface, SEM images of three fields of view were taken at 2000x magnification. From the obtained SEM images, the area ratio of martensite, bainite, and retained austenite, as well as the area ratios of microstructures other than martensite, bainite, and retained austenite (ferrite, cementite), were determined through image analysis. The area ratios of martensite and bainite are calculated by subtracting the area ratio of retained austenite obtained by XRD from the area ratios of martensite, bainite, and retained austenite obtained through image analysis. The average value of the three fields of view is taken as the area ratio of the microstructure.

[0110] Average grain size of the original austenite grains: less than 10 μm

[0111] By refining the grain size, the crack propagation path is made more complex, thereby improving toughness and flexural strength. Furthermore, grain refinement strengthens the grains, thus increasing yield strength. To achieve these effects, the average grain size of the original austenite grains needs to be 10 μm or less. Preferably, the average grain size of the original austenite grains is 9 μm or less. While there is no particular limitation on the lower limit of the average grain size of the original austenite grains, from a manufacturing perspective, it is preferable to set it to 1 μm or more.

[0112] Here, the average grain size of the original austenite grains was determined as follows. After grinding the thickness section of each steel plate parallel to the rolling direction, it was etched with picrol. SEM images of the microstructure at 1 / 4t of the plate thickness were taken at 2000x magnification, capturing three fields of view. The grain size of each original austenite grain was determined by image analysis from the obtained microstructure images, and the average value of the three fields of view was taken as the average grain size of the original austenite grains.

[0113] Bode concentration at the original austenite grain boundaries: ≥0.10% by mass.

[0114] Borosilicate (B) strengthens grain boundaries by segregating at the original austenite grain boundaries, thereby improving toughness and flexural strength. This effect is achieved when the B concentration at the original austenite grain boundaries is 0.10% or more by mass. Preferably, the B concentration at the original austenite grain boundaries is 0.15% or more by mass, more preferably 0.20% or more. Although no upper limit is set for the B concentration at the original austenite grain boundaries, it is preferable to be less than 20% in order to properly prevent the precipitation of hard borocarbons at the grain boundaries and further improve toughness.

[0115] Here, the boron concentration at the proto-austenite grain boundaries was determined as follows. Needle-shaped samples were prepared from the region containing the proto-austenite grain boundaries using SEM-FIB (Focused Ion Beam). The obtained needle-shaped samples were then analyzed using a 3DAP (3Dimensional Atom Probe, 3DAP) apparatus (LEAP4000XSi, AMETEK). Measurements were performed in laser mode. The boron concentration at the proto-austenite grain boundaries was determined based on the number of boron ions detected at the proto-austenite grain boundaries and the number of other ions.

[0116] C-rich areas

[0117] Strengthening martensitic grain boundaries and the parent phase containing them through carbon enrichment can improve bending properties and yield ratio. It should be noted that in this specification, "martensitic grain boundary" includes proto-austenitic grain boundaries, blocky grain boundaries, and packet grain boundaries present in martensite and bainite. Figure 1 The image shows an example of a C-rich region. Figure 1 (A) represents the observations of C-enriched regions present in bulk grain boundaries and lath bundle grain boundaries. Figure 1 (B) is a graph showing the observations of C-rich regions present at the original austenite grain boundaries. Figure 1 In (A) and (B), the left-hand image shows an example of observation results using a Scanning Transmission Electron Microscope (STEM), revealing the presence of martensite grain boundaries in the center of the image. The right-hand image shows an example of STEM observation results regarding carbon enrichment. These images demonstrate that carbon-rich regions exist along the martensite grain boundaries, extending to the parent material containing the martensite grain boundaries.

[0118] The carbon concentration in the carbon-rich region is more than 4.0 times that of the carbon content in the steel.

[0119] In the C-enriched region, sufficient grain boundary strength can be obtained by enriching C to more than 4.0 times the C content in the steel. That is, the C concentration in the C-enriched region satisfies the following equation (2).

[0120] C concentration (mass%) in C-rich regions / C content (mass%) in steel ≥ 4.0…(2)

[0121] The C concentration in the C-enriched region is preferably more than 4.5 times the C content in the steel. There is no specific upper limit for the C concentration in the C-enriched region, but to appropriately prevent the precipitation of cementite and the reduction of the C concentration in the solid solution, a C concentration of 6% or less is preferred.

[0122] C-enriched region: an enrichment width of 3 nm to 100 nm in a direction orthogonal to the martensite grain boundary.

[0123] like Figure 1 As shown, not only the martensitic grain boundaries, but also the parent phase sandwiched between them is strengthened by C enrichment, thereby improving bending properties and yield ratio. Therefore, the C-enriched region is made to exist with an enrichment width of 3 nm to 100 nm in a direction orthogonal to the martensitic grain boundaries. When the enrichment width of the C-enriched region is less than 3 nm, the above effect is small. On the other hand, when the width of the C-enriched region exceeds 100 nm, C cannot be sufficiently enriched at and near the grain boundaries. The width of the C-enriched region is preferably 3.5 nm or more. Furthermore, the width of the C-enriched region is preferably 80 nm or less.

[0124] C-rich regions: lengths exceeding 100 nm in directions parallel to martensite grain boundaries.

[0125] To achieve excellent bending and yield ratio, it is important to strengthen the martensite grain boundaries into a network structure through carbon segregation. Therefore, the carbon-enriched regions should have a length of 100 nm or more in the direction parallel to the martensite grain boundaries. When the carbon-enriched regions are less than 100 nm, failure and yielding occur from the gaps in the carbon-enriched regions. Preferably, the carbon-enriched regions should have 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 carbon-enriched regions along the martensite grain boundaries; the carbon-enriched regions can exist in a manner that covers the entire length of the martensite grain boundaries.

[0126] Here, the C concentration, enrichment width, and length of the C-enriched region were determined as follows. Thin film samples containing martensitic grain boundaries were prepared using SEM-FIB, and surface analysis of C was performed using STEM and Energy Dispersive X-ray Spectroscopy (EDS). The analysis was conducted using an analytical transmission electron microscope, Talos F200X (FEI manufactured). The thin film sample was tilted with the martensitic grain boundaries parallel to the electron beam, and surface analysis was performed on a 200 × 500 nm region. The analysis length in the direction parallel to the martensitic grain boundaries (along the direction of the martensitic grain boundaries) was set to 500 nm. Surface analysis data were accumulated in the direction parallel to the martensitic grain boundaries, and a line profile with a length of 200 nm was obtained in the direction orthogonal to the martensitic grain boundaries. In the C concentration line profile, half the maximum value of the line profile was calculated, and the width of the line profile that is greater than half the maximum value was taken as the enrichment width of the C-enriched region. For this enrichment width, quantitative EDS analysis was performed to quantify the C concentration in the C-enriched region. Additionally, in the surface analysis of C, the length of the C-enriched region was measured in a direction parallel to the martensite grain boundary, and this length was taken as the length of the C-enriched region along the martensite grain boundary.

[0127] According to the present invention, a high-strength steel plate with a tensile strength of 1180 MPa or higher can be provided. Preferably, the tensile strength of the high-strength steel plate is 1250 MPa or higher.

[0128] The high-strength steel sheet described above may have a coating on at least one side. Preferably, the coating is a hot-dip galvanized layer, an alloyed hot-dip galvanized layer, or an electro-galvanized layer. The composition of the coating is not particularly limited and may be a known composition.

[0129] The composition of the hot-dip galvanized layer is not particularly limited, as long as it is a general composition. In one example, the coating has the following composition: containing Fe: less than 20% by mass, Al: 0.001% to 1.0% by mass, and a total of 0% to 3.5% by mass of one or more elements selected from Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, and REM, with the remainder consisting of Zn and unavoidable impurities. In the case of a hot-dip galvanized layer, in one example, the Fe content in the coating is less than 7% by mass; in the case of an alloyed hot-dip galvanized layer, in one example, the Fe content in the coating is 7% to 15% by mass, more preferably 8% to 13% by mass.

[0130] There is no particular limitation on the coating amount, but it is preferred to set the coating amount on each single side of the high-strength steel plate to be 20-80 g / m². 2 In one example, the coating is formed on both the front and back sides of a high-strength steel sheet.

[0131] Next, the manufacturing method of high-strength steel plates will be explained.

[0132] First, a steel slab with the above-mentioned composition is manufactured. First, the steel billet is melted to produce molten steel with the above-mentioned composition. The melting method is not particularly limited; known melting methods such as converter melting and electric furnace melting are suitable. The resulting molten steel is solidified to manufacture a steel slab. The method for manufacturing the steel slab from the molten steel is not particularly limited; continuous casting, ingot casting, or thin slab casting can be used. The steel slab can be hot-rolled after temporary cooling and reheating, or the cast steel slab can be continuously hot-rolled without cooling to room temperature. Considering the rolling load and the formation of oxide scale, the slab heating temperature is preferably set to 1100°C or higher, and more preferably 1300°C or lower. The slab heating method is not particularly limited; for example, it can be heated in a heating furnace using conventional methods.

[0133] [Hot rolling process]

[0134] Next, the heated steel slab is hot-rolled to produce a hot-rolled sheet. There are no particular restrictions on the hot rolling process; conventional methods can be used. There are also no particular restrictions on cooling after hot rolling; the sheet is cooled to the winding temperature. Then, the hot-rolled sheet is wound into coils. The winding temperature is preferably set to 400°C or higher. This is because if the winding temperature is above 400°C, the strength of the hot-rolled sheet will not increase, making winding easier. A winding temperature of 550°C or higher is more preferred. Furthermore, to appropriately prevent the formation of a thick oxide scale and further improve the yield, the winding temperature is preferably set to 750°C or lower. It should be noted that the hot-rolled sheet can also be heat-treated for softening purposes before pickling.

[0135] [Pickling process]

[0136] Optionally, the oxide scale on the hot-rolled sheet wound into a coil is removed. There are no particular limitations on the method for removing the oxide scale; however, to completely remove it, it is preferable to perform pickling while unwinding the hot-rolled coil. There are no particular limitations on the pickling method; conventional methods are acceptable.

[0137] [Cold rolling process]

[0138] Cold-rolled sheets are produced by appropriately cleaning hot-rolled sheets, after the oxide scale has been optionally removed, and then cold-rolling them. There are no particular limitations on the cold-rolling method; conventional methods are acceptable.

[0139] [Annealing process]

[0140] Next, the following annealing process is performed: the cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and held for 10 seconds or higher; then, it is heated to a second heating temperature of 1000°C or higher and 1200°C or lower at an average heating rate of 50°C / s or higher; and within 5 seconds after reaching the second heating temperature, it is cooled to 500°C or lower at a cooling rate of 50°C / s or higher.

[0141] First heating temperature above 850℃ and below 920℃

[0142] Next, the cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and held for at least 10 seconds. Annealing is performed at the first heating temperature within the austenite single-phase domain to obtain a microstructure consisting mainly of martensite and bainite. When the first heating temperature is below 850°C, ferrite is formed, resulting in reduced strength. On the other hand, when the first heating temperature exceeds 920°C, the austenite grain size exceeds 10 μm, making it impossible to refine the grain size in subsequent processes, thus reducing flexibility, toughness, and yield ratio. The first heating temperature is preferably 860°C or higher. Furthermore, the first heating temperature is preferably 900°C or lower.

[0143] Holding time at the first heating temperature: 10 seconds or more

[0144] The holding time at the first heating temperature is set to 10 s or more. By holding the first heating temperature for 10 s or more, the growth of austenite grain size is balanced with the pinning based on Nb carbides or the growth inhibition based on solid solution. When the holding time is less than 10 s, the austenite grains cannot exhibit the pinning effect caused by Nb carbides or the growth inhibition effect caused by solid solution during subsequent rapid heating, and the original austenite grain size exceeds 10 μm. There is no particular upper limit to the holding time at the first heating temperature, but from a productivity point of view, the holding time at the first heating temperature is preferably set to 60 s or less. The holding time at the first heating temperature is preferably 20 s or more.

[0145] Second heating temperature above 1000℃ and below 1200℃

[0146] After holding at the first heating temperature, annealing is performed at a high temperature while maintaining austenite grain boundaries below 10 μm, allowing boron (B) to segregate sufficiently at the grain boundaries. When the second heating temperature is below 1000°C, B diffusion is slow, and grain boundary segregation is insufficient. When the second heating temperature exceeds 1200°C, austenite grain growth is rapid, and the austenite grain size exceeds 10 μm. The second heating temperature is preferably set to 1020°C or higher. Alternatively, the second heating temperature is preferably set to 1150°C or lower.

[0147] Average heating rate: 50℃ / s or higher

[0148] 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. There is no particular upper limit to the average heating rate from the first heating temperature to the second heating temperature, but excessively rapid heating is difficult to control, so it is preferably set to 120°C / s or less. The average heating rate from the first heating temperature to the second heating temperature is preferably 80°C / s or more.

[0149] Within 5 seconds of reaching the second heating temperature, cool to below 500°C at an average cooling rate of 50°C / s or higher.

[0150] Upon reaching the second heating temperature, there is no need to hold it. Rapid cooling begins within 5 seconds of reaching the second heating temperature, with an average cooling rate of 50°C / s or higher, rapidly cooling to below 500°C. This yields a steel microstructure with austenite grain size of less than 10 μm and grain boundary segregation of 0.1% or higher. If the second heating temperature is held, grain growth begins rapidly, and cooling commences immediately upon reaching the second heating temperature.

[0151] Average cooling rate: 50℃ / s or higher

[0152] During cooling after reaching the second heating temperature, the average cooling rate from the second heating temperature to below 500°C is set to 50°C / s or higher. When the average cooling rate from the second heating temperature to below 500°C is less than 50°C / s, grain growth occurs during cooling. There is no particular upper limit to the average cooling rate from the second heating temperature to below 500°C, but for ease of control, it is preferably set to 120°C / s or lower. The average cooling rate from the second heating temperature to below 500°C is preferably set to 80°C / s or higher.

[0153] Cooling stop temperature: below 500℃

[0154] In addition, to suppress the ferrite phase transformation, rapid cooling is performed to a cooling stop temperature below 500°C. The cooling stop temperature is preferably set to below 450°C. There is no particular limitation on the lower limit of the cooling stop temperature, but it is preferably set to above 100°C.

[0155] Alternatively, the plating process can be performed after the annealing process and before the reheating process, that is, plating is applied to at least one side of the high-strength steel sheet to obtain a high-strength plated steel sheet. Alternatively, the high-strength plated steel sheet can be heat-treated after the plating process to alloy the plating layer of the high-strength plated steel sheet to obtain an alloyed plated steel sheet.

[0156] A rolling process that involves rolling an elongation of 0.5% or more after annealing.

[0157] Following the aforementioned annealing process, a rolling process is performed to roll the cold-rolled sheet with an elongation of 0.5% or more to obtain a second cold-rolled sheet. The cold-rolled sheet obtained up to this point contains a large number of movable dislocations. Through this rolling process, the movable dislocations aggregate at grain boundaries and become intertwined, forming immobile dislocations. The effect is minimal when the elongation is less than 0.5%. The elongation in the rolling process is preferably set to 0.6% or more. No specific upper limit is set for the elongation in the rolling process; however, to further reduce the load on the equipment, it is preferably, for example, 2% or less.

[0158] Following the aforementioned rolling process, the second cold-rolled sheet is subjected to a reheating process at a reheating temperature of 70°C or higher and 200°C or lower for at least 600 seconds.

[0159] Following the aforementioned rolling process, carbon segregates or clusters form on dislocations near grain boundaries, thus the second cold-rolled sheet is tempered at a low temperature. When the reheating temperature is below 70°C, carbon diffusion is slow, and carbon does not accumulate to a sufficient amount near the grain boundaries. On the other hand, when the reheating temperature exceeds 200°C, tempering is excessive, and cementite precipitates. The cementite precipitated at the grain boundaries becomes a fracture initiation point, and the carbon concentration in the parent phase surrounding the cementite decreases, thus reducing flexibility and toughness. The reheating temperature is preferably set to 90°C or higher. Alternatively, the reheating temperature is preferably set to 190°C or lower.

[0160] Holding time at reheat temperature: 600s or more

[0161] When the holding time at the reheating temperature is less than 600 s, C diffusion is slow, and a sufficient amount of C enrichment cannot be obtained. There is no particular upper limit to the holding time at the reheating temperature, but to prevent the precipitation of cementite, it is preferably 43200 s (0.5 days) or less. The holding time at the reheating temperature is preferably 800 s or more.

[0162] Furthermore, without the aforementioned rolling process and upon reheating, carbon segregates at the grain boundaries, improving toughness. However, the enrichment width is narrow, and the area outside the grain boundaries is not strengthened, resulting in poor bending properties. Additionally, dislocations remain mobile dislocations, leading to poor YR (yellow-red) properties.

[0163] It should be noted that manufacturing conditions other than those mentioned above can be based on commonly used methods.

[0164] [part]

[0165] This invention provides a component made at least in part of the aforementioned high-strength steel sheet or high-strength clad steel sheet. In one example, the aforementioned high-strength steel sheet or high-strength clad steel sheet can be formed into a target shape by pressing to manufacture an automotive component. It should be noted that the automotive component may also include steel sheets other than the high-strength steel sheet or high-strength clad steel sheet of this embodiment as raw materials. According to this embodiment, a high-strength steel sheet with a TS of 1180 MPa or higher, possessing both bending strength, toughness, and a high yield ratio, can be provided. Therefore, it is suitable as an automotive component that contributes to the lightweighting of the vehicle body. This high-strength steel sheet or high-strength clad steel sheet can be suitably used in automotive components, particularly in all components used as frame structural members or reinforcing members.

[0166] Example

[0167] Steel with the composition shown in Table 1, the remainder consisting of Fe and unavoidable impurities, was smelted in a converter to produce steel slabs. The resulting slabs were then reheated and hot-rolled to obtain hot-rolled coils. Next, while uncoiling the hot-rolled coils, pickling was performed, followed by cold rolling. The thickness of the hot-rolled sheet was 3.0 mm, and the thickness of the cold-rolled sheet was 1.2 mm. Annealing was carried out using a continuous hot-dip galvanizing production line under the conditions shown in Table 2 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 sheets were immersed in a 460°C galvanizing bath at a concentration of 35 g / m² per side. 2 The coating adhesion amount. Alloyed hot-dip galvanized steel sheets are adjusted to 45g / m² per single side. 2 After the coating is applied, an alloying treatment is performed at 520°C for 40 seconds to manufacture the steel sheet. The resulting steel sheet is then rolled and reheated under the conditions shown in Table 2.

[0168] For the obtained steel sheet, the following parameters were evaluated according to the method described above: the combined area ratio of martensite and bainite, the original austenite grain size, the Brønsted boron concentration at the original austenite grain boundaries, the C concentration (mass%) of the C-enriched region at the martensite grain boundaries / the C content (mass%) in the steel, the enrichment width of the C-enriched region, and the length of the C-enriched region along the martensite grain boundaries. Furthermore, tensile strength, yield ratio, toughness, and flexural properties were evaluated according to the method described later. The results are shown in Table 3.

[0169] [Tension Test]

[0170] The obtained steel plates were subjected to tensile tests according to JIS Z 2241. Using the direction orthogonal to the rolling direction as the length direction, JIS No. 5 tensile test specimens were collected and tensile tests were performed to determine the tensile strength (TS) and yield strength (YS). If the tensile strength TS is 1180 MPa or higher, it is considered to have good tensile strength. Furthermore, if the ratio of yield strength to tensile strength, YR = YS / TS, is 0.80 or higher, it is considered to have a high yield ratio.

[0171] [The Charpy Experiment]

[0172] Charpy impact testing was conducted according to JIS Z 2242. Test pieces were collected from the obtained steel plates, with a 90° V-shaped cut perpendicular to the rolling direction, 10 mm wide, 55 mm long, and a 2 mm depth at the center of the length. Charpy impact testing was then performed within the test temperature range of -120 to +120 °C. The transition curve was derived from the obtained brittle fracture surface area, and the temperature at which the brittle fracture surface area reached 50% was defined as the brittle-ductile transition temperature. It should be noted that a brittle-ductile transition temperature below -40 °C obtained from the Charpy test is considered to have good toughness. In the table, toughness is indicated as "excellent" for brittle-ductile transition temperatures below -40 °C, and as "poor" for brittle-ductile transition temperatures above -40 °C.

[0173] [Bending Test]

[0174] The bending test was conducted according to JIS Z 2248. From the obtained steel sheet, a strip-shaped test piece with a width of 30 mm and a length of 100 mm was collected, with the bending test axis parallel to the rolling direction of the steel sheet. Then, a 90°V bending test was performed under the conditions of an indentation load of 100 kN and a holding time of 5 seconds. It should be noted that bending performance was evaluated based on the pass rate of the bending test. Five samples were subjected to bending tests at the maximum R (e.g., a bending radius of 7.0 mm in the case of a 1.2 mm sheet thickness) where the value R / t obtained by dividing the bending radius (R) by the sheet thickness (t) was 5 or less (e.g., bending radius of 7.0 mm in the case of a 1.2 mm sheet thickness). Then, the presence or absence of cracks at the ridge of the bending apex was evaluated. Only when none of the five samples cracked, i.e., the pass rate was 100%, was the bending performance considered good. In the table, bending performance is indicated as "excellent" only when the pass rate is 100%, and "poor" for other cases. Here, the presence or absence of cracks is evaluated by measuring the edge of the bend apex using a digital microscope (RH-2000: manufactured by HIROX Co., Ltd.) at 40x magnification.

[0175]

[0176] Table 2

[0177]

[0178] The underlined part indicates outside the suitable scope of the invention.

[0179]

[0180] According to Table 3, in the present invention example, the TS is 1180 MPa or more, the yield ratio is 0.80 or more, and the bending and toughness are excellent. On the other hand, in the comparative example, one or more of the following are poor: TS, yield ratio, bending and toughness.

Claims

1. A high-strength steel plate having the following composition: by mass % containing C: 0.10%~0.30%, Si: 0.20%~1.20%, Mn: 2.5%~4.0%, P: less than 0.050%, S: less than 0.020%, Al: less than 0.10%, N: less than 0.01%, Ti: less than 0.100%, Nb: 0.002%~0.050%, and B: 0.0005%~0.0050%, the remainder consisting of Fe and unavoidable impurities, and satisfying the following formula (1). Furthermore, the combined area ratio of martensite and bainite is over 95%. The average grain size of the original austenite grains is less than 10 μm. The boron concentration at the original austenite grain boundaries is above 0.10% by mass. There are C-rich regions along the martensite grain boundaries. The C concentration in the C-enriched region, expressed as a percentage by mass, is more than 4.0 times the C content in the steel, expressed as a percentage by mass. It has an enrichment width of 3 nm to 100 nm in a direction orthogonal to the martensite grain boundary, and a length of more than 100 nm in a direction parallel to the martensite grain boundary. ([%N] / 14) / ([%Ti] / 47.9)<1.0…(1) In equation (1), [%N] and [%Ti] represent the mass percentage of N and Ti in the steel, respectively.

2. The high-strength steel plate according to claim 1, wherein, The composition further contains, by mass%, at least one element selected from the following: V: ≤0.100%, Mo: ≤0.500%, Cr: ≤1.00%, Cu: ≤1.00%, Ni: ≤0.50%, Sb: ≤0.200%, Sn: ≤0.200%, Ta: ≤0.200%, W: ≤0.400%, Zr: ≤0.0200%, Ca: ≤0.0200%, Mg: ≤0.0200%, Co: ≤0.020%, REM: ≤0.0200%, Te: ≤0.020%, Hf: ≤0.10%, and Bi: ≤0.200%.

3. A high-strength galvanized steel sheet having a galvanized layer on at least one side of the high-strength steel sheet as described in claim 1 or 2.

4. A method for manufacturing a high-strength steel plate, comprising the following steps: A hot-rolled plate is produced by hot rolling a steel slab having the composition of claim 1 or 2. The hot-rolled sheet is cold-rolled to produce a cold-rolled sheet. The following annealing process is performed: the cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and held for at least 10 seconds; then, it is heated to a second heating temperature of 1000°C or higher and 1200°C or lower at an average heating rate of 50°C / s or higher; and then, within 5 seconds of reaching the second heating temperature, it is cooled to below 500°C at an average cooling rate of 50°C / s or higher. After the annealing process, a rolling process is performed to roll the cold-rolled sheet to an elongation of 0.5% or more to obtain a second cold-rolled sheet. After the rolling process, a reheating process is performed to hold the second cold-rolled sheet at a reheating temperature of 70°C or higher and 200°C or lower for more than 600 seconds to obtain a high-strength steel sheet.

5. A method for manufacturing a high-strength coated steel sheet, comprising a coating process, after the annealing process described in claim 4 and before the reheating process, a coating process of performing a coating treatment on at least one side of the high-strength steel sheet to obtain the high-strength coated steel sheet.

6. A component, at least a portion of which is made of the high-strength steel plate as described in claim 1 or 2.

7. A component, at least a portion of which is made of the high-strength plated steel sheet as described in claim 3.

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