High-strength steel sheet and method for manufacturing the same

Abstract

This invention provides a high-strength steel plate with a tensile strength of 980 MPa or higher and excellent resistance to hydrogen embrittlement in the bending section, as well as a method for manufacturing the same. Specifically, regarding the microstructure at the 1 / 4 thickness position and the outermost layer, the (ferrite + bainite) content is 0–20% by area, the residual γ content is 4–20% by volume, the tempered martensite content is 70% or higher by area, and the fresh martensite content is 0–20% by area. The thickness section along the rolling direction is evaluated using SEM-EBSD, and the texture of the tempered martensite is represented using the Bunge method. In the orientation distribution function shown in the Eulerian space of 0≤φ1≤90°, φ2=45°, 0≤Φ≤90°, for the 1 / 4 position of the plate thickness, the maximum ODF of the {111} grain, expressed as the maximum intensity at φ1=5×n°, φ2=45°, Φ=55° (n is an integer from 0 to 18), is greater than 2.0. For the surface layer of the plate thickness, the maximum ODF of the {100} grain, expressed as the maximum intensity at φ1=5×n°, φ2=45°, Φ=0° (n is an integer from 0 to 18), is greater than 1.0.

Description

Technical Field

[0001] This invention relates to high-strength steel sheets with a tensile strength (TS) of 980 MPa or higher, primarily suitable for structural and electrical components of automobile bodies, and methods for manufacturing the same. Background Technology

[0002] In recent years, with increasing attention to global environmental issues, CO2 emission regulations have become increasingly stringent. In the automotive sector, there are requirements to improve fuel efficiency through lightweight vehicle bodies and to ensure crash safety with the large batteries used in electric vehicles. Therefore, for steel used in automobiles, there is a strong push to adopt high-strength materials that achieve high body rigidity and crash safety even with thinner walls and lighter weight. Many steel manufacturers are expanding their product lineups of high-strength materials with tensile strengths of 980 MPa or even higher for automotive applications.

[0003] Generally, known methods for strengthening steel include solid solution strengthening, precipitation strengthening, and dislocation strengthening (phase transformation strengthening). For high-strength steel sheets with tensile strength exceeding 980 MPa, phase transformation strengthening is needed to increase the proportion of hard phases such as martensite and bainite. However, when the steel sheet contains more of these hard phases, its machinability decreases. Consequently, it cannot be used in components requiring complex machining, resulting in limited applicability. In recent years, the development of automotive steel materials has focused on achieving both high strength and high machinability while maintaining hard phases.

[0004] In high-strength steel sheets utilizing this hard phase, methods using tempered martensite are known to improve mechanical properties. For example, Patent Documents 1 and 2 disclose a technique in which, after cooling to a temperature below the Ms point during the final annealing process, the sheet is reheated to temper the martensite generated during cooling, thereby obtaining high tensile flangeability. Patent Document 3 discloses a method of tempering the steel sheet after plating. Furthermore, Patent Document 4 discloses a method of manufacturing a steel sheet with good bending properties by cooling the steel sheet to below 200°C after plating, and then tempering it in a temperature range of 100°C to 600°C.

[0005] Furthermore, by further refining these technologies, while maintaining a hard, tempered martensite as the main component, the austenitic phase, which provides ductility, is utilized, thus establishing a technology that balances high strength and high ductility in steel sheets, as shown in Patent Document 5. Here, in the final annealing process, a martensitic and austenitic microstructure is formed by cooling to a temperature below the martensitic transformation initiation temperature and above the martensitic transformation end temperature. Then, to stabilize the austenite and temper the martensite, a reheating and holding process is performed.

[0006] Existing technical documents

[0007] Patent documents

[0008] Patent Document 1: Japanese Patent No. 5463685

[0009] Patent Document 2: International Publication No. 2009 / 054539

[0010] Patent Document 3: Japanese Patent Application Publication No. 06-108152

[0011] Patent Document 4: Japanese Patent Application Publication No. 2017-48412

[0012] Patent Document 5: Japanese Patent No. 6787525

[0013] Patent Document 6: Japanese Patent No. 6705561 Summary of the Invention

[0014] As mentioned above, the decrease in ductility due to increased material strength has been a long-standing challenge in this field. Furthermore, with increased material strength, not only do the overall elongation and formability (such as bendability) deteriorate, but cracks also form on the surface of the steel sheet during bending. Moreover, starting from these microcracks, hydrogen embrittlement can lead to fracture. The inventors have discovered that bending fracture becomes particularly pronounced under harsh forming conditions, such as bending angles less than 90°, or under harsh hydrogen intrusion environments, such as those with a pH less than 1.0. To address this problem, Patent Document 4 discloses an invention that achieves good elongation characteristics by forming a surface decarburized layer, illustrating the effect of the soft ferrite in the surface layer on improving bending formability. However, the formation of the soft ferrite in the surface layer not only reduces the strength level of the base material but also promotes the formation of microcracks and voids during processing by increasing the interface between ferrite and martensite, which have different mechanical properties. Particularly in materials with tensile strength exceeding 980 MPa, the problem of fracture accompanied by hydrogen embrittlement exists.

[0015] The inventors have studied a manufacturing method for improving the balance of strength and ductility and resistance to hydrogen embrittlement of a high-strength steel plate with tempered martensite as the main component (70% or more), the remaining part mainly composed of austenite, and a tensile strength (TS) of 980 MPa or more.

[0016] First, following publicly available technical methods, attempts were made to adjust annealing parameters in the final annealing process, including annealing temperature / time, cooling rate, cooling stop temperature / time, and reheating temperature / time. However, essentially, there is a trade-off between the strength (TS in tensile tests) and ductility (total elongation) of the steel sheet, and no significant improvement in the strength-ductility balance was observed. Next, the annealing process was fundamentally re-examined, and an investigation was conducted, as shown in Patent Document 6, to improve properties by subjecting the steel to an additional heat treatment of Ac1 to Ac1+130°C after cold rolling. However, when applying the known method, the added heat treatment resulted in coarsening of the steel microstructure, leading to a decrease in strength, and a tendency for a worsening of the strength-ductility balance was observed.

[0017] The present invention was made in view of the above circumstances. Specifically, its object is to provide a high-strength steel sheet and a method thereof that maintains a strength strength (TS) of 980 MPa or higher, does not utilize a softening layer on the surface, exhibits excellent resistance to hydrogen embrittlement in the bending section, and improves the balance between strength and ductility. Here, the aforementioned bending section refers to the section subjected to particularly stringent forming conditions, such as bending treatment with a bending angle exceeding 30 degrees.

[0018] The inventors have conducted in-depth research to solve the above-mentioned problems.

[0019] The inventors also focused on the upstream manufacturing process, and mainly re-examined the heat treatment method, rather than simply adjusting the temperature and time.

[0020] First, the experiment that led to the development of this invention will be described. The inventors conducted the following experiment to verify whether a fundamental change in the heat treatment mode could improve the balance between strength and ductility, as well as resistance to hydrogen embrittlement.

[0021] <Experiment 1>

[0022] The steel billet, having the chemical composition shown by symbol AA in Table 1 and consisting of Fe and unavoidable impurities, was smelted and rolled into a billet. It was then hot-rolled at 1200°C for a total of 5 passes at an exit-side rolling temperature of 900°C, followed by air cooling to 400°C and coiling. Next, it was uniformly ground from 3.4 mm thickness to 3.0 mm across both sides, then homogenized in a nitrogen atmosphere at 500°C for 3600 minutes, and cold-rolled to 1.4 mm. Following this, it underwent cold rolling in a nitrogen atmosphere at 500°C–850°C for 60 minutes, then reheated, cooled, and pickled to remove any slight oxide scale on the surface. Finally, it was homogenized in a nitrogen atmosphere at 880°C for 60 seconds, gas-cooled to 150°C, reheated to 300°C for 45 seconds and held, and then gas-cooled to room temperature. After pickling the obtained steel plate, cut out JIS No. 5 test pieces (marking distance 50mm, parallel part width 25mm) in a direction perpendicular to the rolling direction, and conduct tensile tests according to JIS Z 2241.

[0023]

[0024] The results are shown in Table 2. Without post-cold rolling heat treatment (symbol 1), the elongation was 9.3%, compared to a maximum elongation of 10.3% with post-cold rolling heat treatment. Furthermore, the effect of heat treatment temperature was clearly significant, with the increase in elongation occurring only in the temperature range above 600°C and below 800°C (below point A1). Regarding strength, although a slightly lower tendency was observed at 850°C, no significant differences were observed under other conditions.

[0025]

[0026] In addition, 50L×100C samples were punched from the obtained steel sheet with a 15% gap, and then subjected to a 60° C-bending process to form a V-shape, resulting in a bending R / t = 4, to prepare test pieces. Next, the opening angle of the bending section of the test piece was varied, and the test piece was constrained under a residual stress of up to 1.2×TS applied to its surface. The constrained test piece was then placed in hydrochloric acid at a liquid temperature of 25°C and a pH of 0.8 for 24 hours. Subsequently, the presence of cracks in the central part of the bending process was visually observed. As a result, cracks were observed in all steel sheets used in Experiment 1.

[0027] The inventors further conducted the following experiments.

[0028] <Experiment 2>

[0029] The steel billet, having the chemical composition shown by symbol BB in Table 1 and consisting of Fe and unavoidable impurities, was melted and rolled into a billet. Then, it was hot-rolled at 1250°C for a total of 6 passes at an exit-side rolling temperature of 850°C, cooled to 400°C in air, and coiled to obtain a hot-rolled steel sheet. The hot-rolled steel sheet was then uniformly ground from 3.6 mm thickness to 3.2 mm across both sides, followed by a preheating treatment in a nitrogen atmosphere at various temperatures for 3600 minutes, and then cooled in the atmosphere. It was then pickled and cold-rolled to 1.0 mm to obtain a cold-rolled steel sheet. This cold-rolled steel sheet was then heat-treated at 675°C in a nitrogen atmosphere for 60 minutes, cooled, and pickled to remove any slight oxide scale formed on the surface. Next, the cold-rolled steel sheet was homogenized at 880°C for 60 seconds in a nitrogen atmosphere. After the gas cooled to 150°C, it was reheated at 300°C for 45 seconds and held, followed by gas cooling to room temperature. After pickling, JIS No. 5 test pieces (marking distance 50 mm, parallel width 25 mm) were cut in a direction perpendicular to the rolling direction, and tensile tests were performed according to JIS Z 2241 (2022).

[0030] The results are shown in Table 3. Without pre-rolling heat treatment (symbol 1), the elongation of the steel sheet was 9.2%, compared to a maximum elongation of 10.1% achieved with pre-rolling heat treatment. Furthermore, the effect of heat treatment temperature was clearly significant, with the increase in elongation only observed in the temperature range above 400°C and below 800°C. Annealing at 800°C resulted in material hardening and inability to roll, making it impossible to evaluate tensile properties. This is attributed to the transformation of austenite generated during annealing at 800°C into a hard structure during cooling. Regarding strength, although a decrease of approximately 10–20 MPa was observed above 400°C, the difference was not significant.

[0031]

[0032] Furthermore, 50L×100C samples were punched from the obtained thin steel sheet with a 15% gap, and then subjected to a 75° C-bending process to form a V-shape, resulting in a bending R / t = 5, to prepare test pieces. The opening angle of the bending section of the test piece was varied, and the test piece was constrained under a residual stress of up to 1.2×TS applied to its surface. The constrained test piece was then placed in hydrochloric acid at a liquid temperature of 25°C and a pH of 0.8 for 200 hours. The presence of cracks in the central part of the bending process was then visually observed. As a result, cracks were observed in all steel sheets used in Experiment 2.

[0033] Based on the preliminary test results above, the inventors discovered that, in order to improve the balance between strength and ductility, a two-stage heat treatment process, involving heat treatment before cold rolling and heat treatment after cold rolling, can improve the balance between strength and ductility. Further details are provided below, and it was discovered that by combining optimizations such as the addition of Cu, Sb, and Sn, which were not previously found, the resistance to hydrogen embrittlement can be improved, thus completing this invention.

[0034] As a result, the inventors have achieved a high degree of control over the crystal orientation of the steel sheet without utilizing the soft ferrite phase by controlling the microstructure before cold rolling, applying ultra-high rolling pressure during cold rolling, and using slow heating or two-stage annealing (heat treatment) for final annealing. This has resulted in excellent resistance to hydrogen embrittlement and the ability to maintain a tensile strength of 980 MPa while improving the balance between strength and ductility. In other words, the essence of this disclosure is as follows.

[0035] [1] A high-strength steel plate, comprising the following components:

[0036] It contains, by mass%, C: 0.090%–0.300%, Si: 0.40%–2.50%, Mn: 1.8%–4.0%, P: less than 0.100%, S: less than 0.0200%, Al: less than 0.200%, N: less than 0.0200%, and O: less than 0.0100%; and contains at least two of the following: Cu: 0.005%–0.5%, Sn: 0.005%–0.5%, and Sb: 0.001%–0.07%, totaling 0.01%–1.0%; and contains at least one of the following: Ti: less than 0.200%, Nb: less than 0.200%, and V: less than 0.200%, with a content of more than 0.01%; the remainder consists of Fe and unavoidable impurities.

[0037] The microstructure of the steel at the 1 / 4 position of the plate thickness meets the following requirements: the total amount of ferrite and bainite is 0%–20% by area ratio, the amount of retained austenite is 4%–20% by volume ratio, the amount of tempered martensite is more than 70% by area ratio, and the amount of fresh martensite is 0%–20% by area ratio.

[0038] Furthermore, the microstructure of the outermost layer of steel meets the following requirements: the total area ratio of ferrite and bainite is 0% to 20%, the volume ratio of retained austenite is 4% to 20%, the area ratio of tempered martensite is more than 70%, and the area ratio of fresh martensite is 0% to 20%.

[0039] The thickness section of the rolled plate was evaluated using the SEM-EBSD method. The texture of the tempered martensite was found in the orientation distribution function (ODF) in the Eulerian space represented by the Bunge method, where 0°≤φ1≤90°, φ2=45°, and 0°≤Φ≤90°.

[0040] For the 1 / 4 plate thickness position, the maximum ODF of the {111} grain, expressed as the maximum strength at φ1=5×n°, φ2=45°, Φ=55° (n=0, 1, 2, …, 18), is greater than 2.0.

[0041] For the outermost layer, the maximum ODF of the {100} grains, expressed as the maximum intensity at φ1=5×n°, φ2=45°, Φ=0° (n=0, 1, 2, …, 18), is greater than 1.0.

[0042] The tensile strength of the aforementioned high-strength steel plate is above 980 MPa.

[0043] [2] The high-strength steel plate according to [1], wherein the above composition, by mass%, further contains at least one element selected from Ta: less than 0.10%, W: less than 0.10%, B: less than 0.0100%, Cr: less than 1.00%, Mo: less than 1.00%, Co: less than 0.010%, Ni: less than 1.00%, Ca: less than 0.0100%, Mg: less than 0.0100%, REM: less than 0.0100%, Zr: less than 0.100%, Te: less than 0.100%, Hf: less than 0.10%, and Bi: less than 0.200%.

[0044] [3] According to the high-strength steel plate described in [2], the above-mentioned composition also contains, by mass %: Ti: 0.010% to 0.200%, Nb: 0.005% to 0.200% and B: 0.0005% to 0.0100%.

[0045] [4] The high-strength steel plate according to any one of [1] to [3], wherein an electroplated zinc layer is formed on the surface.

[0046] [5] The high-strength steel sheet according to any one of [1] to [3], wherein a hot-dip galvanized layer is formed on the surface.

[0047] [6] The high-strength steel plate according to any one of [1] to [3], wherein an alloyed hot-dip galvanized layer is formed on the surface.

[0048] [7] A method for manufacturing a high-strength steel plate, which is the method for manufacturing a high-strength steel plate as described in any one of [1] to [3],

[0049] After heating the steel billet with the above-mentioned composition to a temperature range of 1080℃~1300℃

[0050] Finish rolling is performed at the finishing temperature within the temperature range of 850℃ to 1000℃.

[0051] After finishing rolling, the temperature is cooled to below 720°C within 2 seconds at a cooling rate of over 70°C per second.

[0052] After winding in a temperature range below Ms+100℃

[0053] Prior to cold rolling, the material undergoes softening annealing for at least 48 hours in a temperature range above 400°C and below Ac1, with or without temperature holding.

[0054] Then, cold rolling is performed, wherein at least one of the first or second passes has a rolling rate of 30% or more, and the total reduction rate is 50% or more. This process yields a cold-rolled sheet.

[0055] The cold-rolled sheet is annealed for at least 60 seconds in a temperature range above 600°C and below the Ac1 point, with or without holding the temperature. Then, it is heated to above the Ac1 point and below 950°C.

[0056] Furthermore, the average cooling rate up to at least 500°C is 10°C / s or higher.

[0057] Cooling to the cooling stop temperature range of 50℃ to 350℃

[0058] Next, heat to a temperature range above the cooling stop temperature but above 250°C and below 600°C, and maintain that temperature range for at least 10 seconds.

[0059] [8] The method for manufacturing high-strength steel plate according to [7] further includes electro-galvanizing.

[0060] [9] The method for manufacturing high-strength steel plate according to [7] further includes hot-dip galvanizing.

[0061]

[10] The method for manufacturing high-strength steel plate according to [7] further includes hot-dip galvanizing, followed by an alloying treatment of galvanizing in a temperature range of 470°C to 600°C.

[0062] According to the present invention, a higher balance of strength and ductility than before can be obtained. As a result, the applicability to automotive body parts is expanded, and it can also be applied to parts requiring complex stamping processes, making a significant contribution to the lightweighting of automotive bodies. In addition, it can improve the resistance to hydrogen embrittlement in processed parts, especially bending parts with particularly harsh forming conditions, which are inherent problems of high-strength materials. Detailed Implementation

[0063] The embodiments of this disclosure will now be described. It should be noted that this disclosure is not limited to the following embodiments. First, an appropriate range of the composition of the steel billet used in the high-strength steel plate of this invention and the reasons for this limitation will be explained. In the following description, unless otherwise specified, "%" indicating the content of the constituent elements of steel means "mass %".

[0064] [C: 0.090%~0.300%]

[0065] Carbon (C) is one of the essential components of steel, and especially in this invention, it is a crucial element affecting the proportions of martensite, ferrite, and retained austenite. When the C content is less than 0.090%, the martensite content decreases, making it difficult to achieve the desired elongation (TS). Furthermore, due to the reduced carbide content and the decrease in {100} grains, cracking during bending tests becomes significant. On the other hand, if the C content exceeds 0.300%, the martensite becomes embrittled, making it difficult to achieve the desired elongation. Therefore, the C content is set to 0.090% to 0.300%. The preferred lower limit is 0.120% or more, more preferably 0.150% or more. The preferred upper limit is 0.280% or less, and even more preferably 0.240% or less.

[0066] [Si: 0.40%~2.50%]

[0067] Si is one of the essential components of steel, especially in this invention. It is an element that affects the hardness of martensite and the percentage of retained austenite because it inhibits carbide formation during continuous annealing and promotes the formation of retained austenite. When the Si content is less than 0.40%, the percentage of retained austenite decreases, making it difficult to achieve the desired El. On the other hand, if the Si content exceeds 2.50%, the carbon concentration in the retained austenite increases excessively, reducing local ductility. Furthermore, bending properties also decrease. Therefore, the Si content is set to 0.40% to 2.50%. The preferred lower limit is 0.60% or more, more preferably 0.80% or more. The preferred upper limit is 2.00% or less, more preferably 1.80% or less.

[0068] [Mn: 1.8%~4.0%]

[0069] Mn is one of the essential components of steel, and especially in this invention, it is a crucial element affecting the martensite content. When the Mn content is less than 1.8%, the martensite content decreases, making it difficult to achieve a temper strength (TS) of 980 MPa or higher. On the other hand, if the Mn content exceeds 4.0%, the tempered martensite content decreases, leading to reduced local ductility. Furthermore, bending properties also decrease. Therefore, the Mn content is set between 1.8% and 4.0%. The preferred lower limit is 2.0% or more, more preferably 2.2% or more. The preferred upper limit is 3.8% or less, more preferably 3.6% or less.

[0070] [P: below 0.100%]

[0071] Phosphorus (P) segregates at the original austenite grain boundaries, causing grain boundary embrittlement and thus reducing the local ductility of the steel sheet, resulting in lower elongation. Additionally, bending resistance is also reduced. Therefore, the P content needs to be 0.100% or less. It should be noted that there is no specific lower limit for the P content, but since P is a solid solution strengthening element that can improve the strength of the steel sheet, it is preferably 0.001% or more. Therefore, the P content is 0.100% or less, preferably 0.070% or less.

[0072] [S: below 0.0200%]

[0073] Sulfur (S), present as a sulfide, reduces the local ductility of the steel sheet, thus decreasing elongation. It also reduces flexibility. Therefore, the S content needs to be set below 0.0200%. It should be noted that while there is no specific lower limit for the S content, from a production technology perspective, it is preferably 0.0001% or higher. Therefore, the S content is 0.0200% or lower, preferably 0.0050% or lower.

[0074] [A1: below 0.200%]

[0075] Adding large amounts of Al will raise the A3 phase transformation point, resulting in a high content of ferrite in the microstructure, thus hindering the achievement of high strength through martensite. Therefore, the Al content needs to be below 0.200%. It should be noted that there is no specific lower limit for the Al content, but since Al can suppress carbide formation and promote the formation of retained austenite during continuous annealing, the Al content is preferably 0.001% or more. A preferred Al content is 0.150% or less.

[0076] [N: below 0.0200%]

[0077] Nitrogen (N), present as a nitride, reduces the local ductility of the steel sheet, thus decreasing elongation. It also reduces flexibility. Therefore, the N content needs to be 0.0200% or less. It should be noted that while there is no specific lower limit for the N content, from a production technology perspective, the N content is preferably 0.0001% or more. Therefore, the N content is 0.0200% or less, and more preferably 0.0100% or less.

[0078] [O: below 0.0100%]

[0079] O, as an oxide, reduces the local ductility of the steel sheet, thus decreasing its elongation. It also reduces its flexibility. Therefore, the O content needs to be below 0.0100%. It should be noted that there is no specific lower limit for the O content, but from a production technology perspective, the O content is preferably 0.0001% or more. Therefore, the O content is 0.0100% or less, preferably 0.0050% or less.

[0080] [At least one of the following: Ti: less than 0.200%, Nb: less than 0.200%, and V: less than 0.200% is 0.01% or more]

[0081] Ti, Nb, and V, by combining with carbon or nitrogen, refine the crystal structure, thereby improving the strength of the material. Therefore, it is necessary to contain at least one of them at 0.01% or more. More preferably, it is 0.015% or more, and even more preferably 0.04% or more.

[0082] It should be noted that when there are two or more types, their total is 0.01% or more.

[0083] Furthermore, if the content of Ti and Nb is 0.200% or less each, large coarse precipitates or inclusions will not be generated, the local ductility of the steel sheet will not be reduced, and therefore the elongation will not be reduced. Additionally, the flexibility will not be reduced. Therefore, the content of Ti and Nb is preferably 0.200% or less each. It should be noted that there is no particular lower limit for the content of Ti and Nb, but since fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet, the content of Ti and Nb is more preferably 0.001% or more each. Therefore, when Ti and Nb are present, their contents are set to 0.200% or less each. More preferably, the content of Ti and Nb is 0.001% or more each. Even more preferably, the content of Ti and Nb is 0.100% or less each.

[0084] If V is 0.200% or less, large coarse precipitates or inclusions will not be generated, the local ductility of the steel sheet will not be reduced, and therefore the elongation will not be reduced. Furthermore, the flexibility will not be reduced. Therefore, the V content is preferably 0.200% or less. It should be noted that there is no particular lower limit for the V content, but since fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet, the V content is more preferably 0.001% or more. Therefore, in the case of V, its content is 0.200% or less. More preferably, the V content is 0.001% or more. Even more preferably, the V content is 0.100% or less.

[0085] [Cu: 0.005% to 0.5%, Sn: 0.005% to 0.5%, Sb: 0.001% to 0.07%]

[0086] If Cu and Sn content is 0.5% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the bendability will not decrease. Therefore, the content of Cu and Sn is preferably 0.5% or less. It should be noted that there is no particular lower limit for the content, but since these are elements that improve hardenability, the content is preferably 0.005% or more. More preferably 0.01% or more. Even more preferably 0.30% or less.

[0087] If the Sb content is 0.07% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the Sb content is preferably 0.07% or less. It should be noted that there is no specific lower limit for the Sb content, but to achieve the effect of suppressing surface decarburization, the content is more preferably 0.001% or more, and even more preferably 0.004% or more.

[0088] [The total percentage of at least two of Cu, Sn, and Sb is 0.01% to 1.0%]

[0089] Cu, Sn, and Sb, as described below, are features of this invention, and the invention must contain at least two of these elements, with a combined content of at least 0.01%. These elements are known to be surface segregation elements, and by adding a certain amount, they can suppress surface decarburization during the annealing process, thereby suppressing microstructure inhomogeneity in the thickness direction and the decrease in base material strength associated with softening of the surface layer. More preferably, the combined content of at least two of Cu, Sn, and Sb is at least 0.015%.

[0090] According to one embodiment of the present invention, a high-strength steel sheet has a composition containing the above-described components, with the remainder consisting of Fe and unavoidable impurities. Furthermore, it is preferable that the high-strength steel sheet according to one embodiment of the present invention has a composition containing the above-described components, with the remainder consisting of Fe and unavoidable impurities. Here, unavoidable impurities include H, Zn, Pb, As, Se, Ge, Sr, and Cs. A total content of 0.100% or less of these impurities is permissible.

[0091] In addition to the above-mentioned composition, the high-strength steel plate of the present invention may further contain, by mass percent, at least one element selected individually or in combination from Ta: less than 0.10%, W: less than 0.10%, B: less than 0.0100%, Cr: less than 1.00%, Mo: less than 1.00%, Co: less than 0.010%, Ni: less than 1.00%, Ca: less than 0.0100%, Mg: less than 0.0100%, REM: less than 0.0100%, Zr: less than 0.100%, Te: less than 0.100%, Hf: less than 0.10%, and Bi: less than 0.200%.

[0092] If Ta and W are each 0.10% or less, large amounts of coarse precipitates or inclusions will not be generated, the local ductility of the steel sheet will not be reduced, and therefore the elongation will not be reduced. Furthermore, the flexibility will not be reduced. Therefore, the content of Ta and W is preferably 0.10% or less each. It should be noted that there is no particular lower limit for the content of Ta and W, but since fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, the strength of the steel sheet is improved. Therefore, the content of Ta and W is more preferably 0.01% or more each. Therefore, when Ta and W are present, their contents are 0.10% or less each. More preferably, the content of Ta and W is 0.01% or more. Even more preferably, the content of Ta and W is 0.08% or less.

[0093] If Cr and Mo are each 1.00% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the content of Cr and Mo is preferably 1.00% or less each. It should be noted that there is no specific lower limit for the content of Cr and Mo, but since they are elements that improve hardenability, the content of Cr and Mo is more preferably 0.01% or more each. Therefore, in the case of containing Cr and Mo, their contents are set to 1.00% or less each. More preferably, the content of Cr and Mo is 0.01% or more. Even more preferably, the content of Cr and Mo is 0.80% or less.

[0094] If the Co content is 0.010% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the Co content is preferably 0.010% or less. It should be noted that there is no specific lower limit for the Co content, but since it is an element that improves hardenability, the Co content is more preferably 0.001% or more. Therefore, in the case of Co content, it is set to 0.010% or less. A more preferred Co content is 0.001% or more. A further preferred Co content is 0.008% or less.

[0095] If the Ni content is 1.00% or less, cracks will not form inside the steel sheet during casting or hot rolling, and the local ductility of the steel sheet will not be reduced, thus the elongation will not be reduced. Furthermore, the flexibility will not be reduced. Therefore, the Ni content is preferably 1.00% or less. It should be noted that there is no particular lower limit for the Ni content, but since Ni is an element that improves hardenability (generally, it is an element that improves corrosion resistance), the Ni content is more preferably 0.001% or more. Therefore, in the case of Ni, its content is set to 1.00% or less, more preferably 0.001% or more. A further preferred Ni content is 0.500% or less.

[0096] If the contents of Ca, Mg, and REM are each 0.0100% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the contents of Ca, Mg, and REM are preferably 0.0100% or less each. It should be noted that there is no particular lower limit for the contents of Ca, Mg, and REM, but since they are elements that spheroidize the shape of nitrides or sulfides and improve the local ductility of the steel sheet, the contents of Ca, Mg, and REM are more preferably 0.0005% or more each. Therefore, in the case where Ca, Mg, and REM are present, their contents are set to 0.0100% or less each. More preferably, the contents of Ca, Mg, and REM are 0.0005% or more. Even more preferably, the contents of Ca, Mg, and REM are 0.0050% or less. It should be noted that REM (rare earth elements) refers to the collective name of 15 elements, including Sc, Y, and lanthanum (La) from atomic number 57 to lutetium (Lu) from atomic number 71. The REM content mentioned here refers to the total content of these elements.

[0097] If Zr and Te are each 0.100% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the content of Zr and Te is preferably 0.100% or less. It should be noted that there is no particular lower limit for the content of Zr and Te, but since they are elements that spheroidize the shape of nitrides or sulfides and improve the local ductility of the steel sheet, the content of Zr and Te is more preferably set to 0.001% or more, respectively. Therefore, in the case of containing Zr and Te, their contents are each set to 0.100% or less. More preferably, the content of Zr and Te is 0.001% or more. Even more preferably, the content of Zr and Te is 0.080% or less.

[0098] If Hf is 0.10% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the Hf content is preferably 0.10% or less. It should be noted that there is no particular lower limit for the Hf content, but since it is an element that spheroidizes nitrides or sulfides and improves the local ductility of the steel sheet, the Hf content is more preferably 0.01% or more. Therefore, in the case of Hf, its content is set to 0.10% or less. A more preferred Hf content is 0.01% or more. A further preferred Hf content is 0.08% or less.

[0099] If the Bi content is 0.200% or less, coarse precipitates or inclusions will not increase, the local ductility of the steel sheet will not decrease, and therefore the elongation will not decrease. Furthermore, the flexibility will not decrease. Therefore, the Bi content is preferably 0.200% or less. It should be noted that there is no particular lower limit for the Bi content, but since it is an element that reduces segregation, the Bi content is more preferably 0.001% or more. Therefore, in the case of Bi, its content is set to 0.200% or less. A more preferred Bi content is 0.001% or more. A further preferred Bi content is 0.100% or less.

[0100] If the content of B is 0.0100% or less, cracks will not be generated inside the steel sheet during casting or hot rolling, and the local ductility of the steel sheet will not be reduced, thus the elongation will not be reduced. Furthermore, the flexibility will not be reduced. Therefore, the content of B is preferably 0.0100% or less. It should be noted that there is no particular lower limit for the content of B, but since it is an element that segregates at the austenite grain boundaries during annealing and improves hardenability, the content of B is more preferably 0.0003% or more. Therefore, in the case of containing B, its content is set to 0.0100% or less. A more preferred content of B is 0.0003% or more. A further preferred content of B is 0.0080% or less.

[0101] Among the aforementioned elements, Ti, Nb, and B are actively utilized in high-strength steel plates due to their relatively low cost. However, they share the common characteristic of delaying recrystallization in steel structures. This effect is particularly pronounced when Ti, Nb, and B are added simultaneously. As described later, this invention is a technique for suppressing austenitic phase transformation from elongated, non-recrystallized structures during final annealing, and therefore exhibits excellent results, especially in steels where Ti, Nb, and B are added simultaneously. The range of Ti, Nb, and B is: Ti: 0.010%–0.200%, Nb: 0.005%–0.200%, B: 0.0005%–0.0100%.

[0102] It should be noted that the Ti, Nb, V, Ta, W, B, Cr, Mo, Ni, Co, Cu, Sn, Sb, Ca, Mg, REM, Zr, Te, Hf and Bi mentioned above are contained as unavoidable impurities when their contents are less than the preferred lower limit values, since they do not impair the effect of the present invention.

[0103] Next, the microstructure will be explained. The locations of the microstructure are defined as the 1 / 4 thickness position and the outermost layer. First, the microstructure at the 1 / 4 thickness position, which is generally considered to have the average characteristics and composition of a steel plate, will be explained.

[0104] [The total ferrite and bainite content is 0%–20%. Field area ratio observed under an electron microscope.]

[0105] Ferrite is a soft microstructure, thus effective in improving workability; however, in the 980 MPa grade high-strength steel sheet, the subject of this invention, it significantly reduces the sheet's strength. The ferrite referred to here can be any of polygonal ferrite, quasi-polygonal ferrite, or granular bainitic ferrite. Furthermore, bainite, like martensite, is an aggregate of lath-like grains, but is softer than martensite. Here, bainite includes upper bainite and lower bainite. For the purpose of achieving high strength, the total content of ferrite and bainite is 20% or less. In a more preferred condition, the total content of ferrite and bainite is 15% or less. Furthermore, a further preferred condition is that the ferrite content is less than 2% and the total content of ferrite and bainite is 10% or less. It should be noted that since these ferrites and bainite may not be included, the lower limit is set to 0% or more.

[0106] The area ratio of ferrite was determined using the following method. First, a section of the steel plate parallel to the rolling direction (L section) was ground and etched with a 3 vol.% nitric acid ethanol solution. Ten fields of view were observed at 2000x magnification at the 1 / 4 thickness position using a SEM (scanning electron microscope). It should be noted that the 1 / 4 thickness position refers to the location corresponding to 1 / 4 of the plate thickness in the depth direction from the steel plate surface. Additionally, the surface of the steel plate in the presence of galvanization refers to the interface between the galvanized layer and the base steel plate. Next, using the obtained microstructure images, the area ratio of the microstructure (ferrite) in the 10 fields of view was calculated using Image-Pro from Media Cybernetics. The average of these 10 field-view area ratios was taken as the "ferrite area ratio". It should be noted that in the above microstructure images, ferrite appears as a gray microstructure (base microstructure). It should be noted that retained austenite and martensite appear as white microstructures.

[0107] [Retained austenite 4%–20%. Volume fraction based on X-ray diffraction]

[0108] Retained austenite contributes to the ductility of the steel sheet due to the TRIP effect and is an essential microstructure in this invention. To achieve a sufficient effect, the lower limit is set to 4% or more. However, if there is an excessive amount of retained austenite, it may transform into a hard martensitic structure when used in molding applications such as automotive parts, potentially impairing processing characteristics. Therefore, the upper limit is set to 20% or less. A preferred lower limit is 6% or more. A preferred upper limit is less than 16%. It should be noted that the volume fraction of retained austenite can be converted to an area fraction.

[0109] The method for determining the volume fraction of retained austenite is as follows: The steel plate is mechanically ground to 1 / 4 of its thickness along the thickness direction (depth direction), and then chemically ground with oxalic acid to create an observation surface. This observation surface is then observed using X-ray diffraction. Using a Co Kα ray source as the incident X-ray, the diffraction peak intensities of fcc iron (austenite) relative to the {200}, {211}, and {220} planes of bcc iron are measured, along with the diffraction peak intensities of the {200}, {220}, and {311} planes. Next, by multiplying by a correction factor, the diffraction peaks of martensite and austenite are averaged separately to eliminate the influence of preferred orientation, and the volume fraction of retained austenite is calculated. Rγ(200)=33.0, Rγ(220)=17.6, Rγ(311)=25.1, Rα(200)=13.2, Rα(211)=27.3, Rα(220)=12.5.

[0110]

[0111] [Tempered martensite content is over 70%. Field area ratio observed under an electron microscope.]

[0112] Tempered martensite is an aggregate of lath-shaped grains characterized by the presence of iron-based carbides within it. Compared to fresh martensite, it exhibits higher ductility and, more importantly, readily yields the high yield strength required to ensure the crash safety of automotive components. To achieve these effects, the area fraction of tempered martensite needs to be 70% or more. More preferably, the area fraction of tempered martensite is 75% or more, and even more preferably 80% or more. It should be noted that there is no particular upper limit to the area fraction of tempered martensite, but to ensure the area fraction of retained austenite, the area fraction of tempered martensite is preferably 94% or less. It should be noted that the area fraction of tempered martensite can be determined by the method described later.

[0113] [Fresh martensite comprises 0%–20%. Field area ratio observed under an electron microscope.]

[0114] Fresh martensite is martensite that does not contain iron-based carbides. While steel sheets containing fresh martensite have high strength, their processing characteristics are poor. Therefore, the fresh martensite area ratio of the steel sheet according to this embodiment is limited to 20% or less. On the other hand, even if the fresh martensite area ratio is 0%, the steel sheet according to this embodiment has sufficient strength. Therefore, the steel sheet according to this embodiment does not need to contain fresh martensite, so the lower limit for the fresh martensite area ratio is 0%. A more preferred range is 10% or less.

[0115] Furthermore, in the steel microstructure of this invention, in addition to the aforementioned ferrite, tempered martensite, bainite, fresh martensite, and retained austenite, even if it contains carbides such as pearlite and cementite, as well as other microstructures known as steel sheet microstructures, as long as their total area fraction is less than 8%, it will not impair the effectiveness of this invention. It should be noted that other steel sheet microstructures (residual microstructures) can be confirmed and determined, for example, by SEM observation.

[0116] The method for determining the area ratios of tempered martensite, fresh martensite, pearlite, and bainite is as follows. Annealed steel sheet samples were cut out, and the section with the thickness parallel to the rolling direction was ground. The sample was then etched with a 3 vol.% nitric acid-ethanol solution. Three fields of view were captured at 1500x magnification at the 1 / 4 thickness position (equivalent to 1 / 4 of the plate thickness in the depth direction from the steel sheet surface). Higher magnification observations could be performed for more detailed confirmation of carbides. It should be noted that the surface of the steel sheet in the presence of galvanization refers to the interface between the galvanized layer and the steel sheet. Using the obtained image data, the area ratio of each microstructure was calculated using Image-Pro from Media Cybernetics. The average area ratio of each microstructure in each field of view was taken as the area ratio of each microstructure. In the above image data, fresh martensite was divided into white or light gray areas, tempered martensite into gray or dark gray areas containing inconsistently oriented carbides, and pearlite into black and white layered structures. Furthermore, gray or dark gray regions containing uniformly oriented carbides are bainite. However, in actual observation, the above classification is sometimes difficult; for example, even in tempered martensite, uniform carbide orientation can sometimes be observed. Therefore, the microstructure containing gray or dark gray regions with uniformly oriented carbides and exhibiting a linearly extending microstructure interface is classified as bainite, while other gray or dark gray regions containing carbides are classified as tempered martensite. Additionally, distinguishing fresh martensite from retained austenite, which also appears in white or light gray regions, is difficult. Therefore, after calculating the area ratio of the white or light gray regions, the volume ratio of retained austenite obtained using the above method is considered as the area ratio and subtracted to obtain the area ratio of fresh martensite.

[0117] Next, the microstructure of the outermost layer will be described. It should be noted that the outermost layer is selected as the region excluding unstable portions such as oxide scale on the very surface. The outermost layer refers to a position 20 μm deep into the plate thickness from the surface of the steel sheet. Similar to the 1 / 4 thickness position, in the presence of a galvanized layer, the interface between the galvanized layer and the base steel sheet is used as the steel sheet surface. The microstructure of the outermost layer is also included within the microstructure fraction defined above as 1 / 4 thickness. However, since the microstructure is uniform throughout the plate thickness, the following description is more preferable.

[0118] A characteristic of this invention is that the microstructure at the outermost layer is not significantly different from that at the 1 / 4-thickness position. Therefore, the values ​​of each of the above-mentioned microstructures are preferably within ±5% of the 1 / 4-thickness position. Since limited field-of-view observations include certain deviations, they will not be completely identical to those at the 1 / 4-thickness position, but it is further preferred that the values ​​are within ±3% of the 1 / 4-thickness position. For example, Patent Document 4 discloses an invention that obtains good elongation properties by forming a surface decarburized layer. The feature of this invention is that it suppresses this surface unstable microstructure that is intentionally formed or previously unavoidable during annealing, distinguishing it from the above-mentioned invention.

[0119] [Tensile strength of 980 MPa or higher]

[0120] In this invention, the tensile strength of the steel plate is set to 980 MPa or higher. However, the strength level is not limited to 980 MPa; high strengths of around 1180 MPa, 1300 MPa, 1470 MPa, and even 1.5 GPa are also considered. Furthermore, the manufacturing method described in this invention is independent of the strength level of the steel plate and can be applied to various materials.

[0121] [The thickness section of the rolled plate was evaluated using the SEM-EBSD method. The texture of the tempered martensite, in the orientation distribution function (ODF) shown in the Eulerian space (0°≤φ1≤90°, φ2=45°, 0°≤Φ≤90°) expressed by the Bunge method, showed that for the 1 / 4 thickness position, the maximum ODF of the {111} grains, expressed as the maximum intensity at φ1=5×n°, φ2=45°, Φ=55° (n=0, 1, 2, ..., 18), was greater than 2.0. For the surface layer of the plate, the maximum ODF of the {100} grains, expressed as the maximum intensity at φ1=5×n°, φ2=45°, Φ=0° (n=0, 1, 2, ..., 18), was greater than 1.0.]

[0122] The texture of the aforementioned bcc-phase grains is characterized by satisfying the above conditions in the orientation distribution function (ODF) expressed in Eulerian space (0°≤φ1≤90°, φ2=45°, 0°≤Φ≤90°) using the Bunge method. It is known that crystal-oriented grains (here referred to as {111} grains) represented by φ1=5×n°, φ2=45°, and Φ=55° are orientation grains that increase local elongation. This crystal orientation is utilized in soft steel sheets to ensure ductility, but in high-strength steel sheets, its effectiveness is not fully understood because the technology for controlling crystal orientation has not yet been fully established.

[0123] ODF was measured using the SEM-EBSD method described above, employing OIM Analysis 8.0 manufactured by TSL Corporation. In the surface layer measurement, the measurement range was defined as 20 μm from the surface to the thickness direction. While the EBSD method was used for ODF measurement, other known methods such as XRD are also acceptable. The inventors investigated the reasons for the improved ductility achieved by heat treatment before cold rolling in <Experiment 2> above, and examined the texture of the resulting steel sheet. The results showed that high ductility was observed when the maximum ODF of the {111} grains, expressed as the maximum strength at 1 / 4 of the sheet thickness (φ1 = 5 × n°, φ2 = 45°, Φ = 55° (n = 0, 1, 2, ..., 18), was 2.0 or higher. More preferably, it was 2.5 or higher.

[0124] Furthermore, the inventors have discovered that by increasing the formation of crystal-oriented grains (referred to herein as {100} grains) with a higher concentration of these grains on the surface of the steel plate, resistance to hydrogen embrittlement can be improved. While the detailed mechanism is not yet fully understood, it is believed that iron is less likely to dissolve in the corrosion reaction of the ND / / {100} grains, represented by the aforementioned Euler angles, during hydrochloric acid immersion, thus inhibiting hydrogen intrusion into the steel. This high resistance to hydrogen embrittlement is manifested when the maximum ODF of the {100} grains, represented by the maximum strength at the outermost layer (φ1=5×n°, φ2=45°, Φ=0° (n=0, 1, 2, ..., 18), is 1.0 or higher. More preferably, it is 1.5 or higher.

[0125] For the aforementioned high-strength steel plates, electro-galvanized layers, hot-dip galvanized layers, or alloyed hot-dip galvanized layers can also be formed. These coatings can be formed using conventional methods.

[0126] Next, the manufacturing method of the high-strength steel plate of the present invention will be described.

[0127] Steel billets with the above-mentioned composition are smelted using conventional refining processes and then produced into steel billets via conventional ingot-rolling or continuous casting. Alternatively, thin steel billets with a thickness of 100 mm or less can be manufactured by direct casting. These steel billets are heated and maintained at 1080°C to 1300°C, then hot-rolled to produce hot-rolled plates through roughing and finishing rolling, and finally coiled. Here, after finishing rolling, the plates are cooled to below 720°C within 2 seconds at a cooling rate of 70°C or higher per second. The thickness of the hot-rolled plate is preferably 0.8 mm to 4.0 mm. The reasons for these limitations on the hot-rolling conditions are as follows.

[0128] [Heating temperature: 1080℃~1300℃]

[0129] Below 1080°C, the precipitates in the steel are not fully dissolved, resulting in a lack of fine microstructure in subsequent heat treatment processes, which is detrimental to achieving high strength. Therefore, the lower limit is 1080°C or higher, and more preferably 1120°C or higher. Furthermore, if the temperature exceeds 1300°C, the steel billet is prone to deformation, increasing the frequency of defects such as dents. Therefore, the upper limit is 1300°C or lower, and more preferably 1280°C or lower.

[0130] [Finishing temperature: 850℃~1000℃]

[0131] For hot rolling, in order to improve the strength-ductility balance of the material through homogenization of the microstructure within the steel sheet, the finishing process needs to end in the austenitic single-phase region; therefore, the finishing temperature is 850°C or higher. Furthermore, if the finishing temperature exceeds 1000°C, the hot-rolled microstructure becomes coarse, and consequently, the microstructure after annealing also becomes coarse, resulting in a decrease in the steel sheet strength. The preferred lower limit of the finishing temperature is 875°C or higher. The preferred upper limit of the finishing temperature is 950°C or lower.

[0132] [Cooling after hot rolling: Cool to below 720°C within 2 seconds at an average cooling rate of over 70°C per second]

[0133] If the temperature is maintained in the range exceeding 720°C for an extended period, the precipitates such as Ti coarsen, resulting in a loss of the strength-enhancing effect. Consequently, even if other manufacturing conditions are adjusted to increase strength, the desired ductility cannot be achieved. Furthermore, the increased grain size before cold rolling leads to a reduction in the {111} grains formed in the final annealed steel sheet. Therefore, cooling should begin within 2 seconds after finishing rolling and be reduced to below 720°C. More preferably, within 1 second. During cooling, from the viewpoint of suppressing precipitate coarsening, the plate temperature should also be as low as possible, thus the average cooling rate up to 720°C should be 70°C or more per second. More preferably, it should be 90°C or more per second. Cooling as referred to here means water cooling, oil cooling, or gas cooling for the purpose of reducing the steel sheet temperature, and the average cooling rate up to 720°C refers to the average cooling rate from the start of cooling until the steel sheet temperature reaches 720°C.

[0134] [Winding temperature: Ms +100℃ or below]

[0135] After finishing rolling, the steel sheet is cooled and then coiled at a temperature below Ms + 100°C to complete the hot rolling process. If the coiling temperature exceeds Ms + 100°C, the steel sheet will produce a large amount of oxide scale and grain boundary oxidation, which not only worsens pickling properties but also makes the hot-rolled microstructure coarser, making it difficult to generate {111} grains that are beneficial for improving local ductility after cold rolling and annealing. Here, Ms is the temperature expressed by the following formula.

[0136] Ms=561-474[C]-33[Mn]-17[Ni]-17[Cr]-21[Mo]

[0137] Here, [C], [Mn], [Ni], [Cr], and [Mo] represent the content (mass%) of each element, and are set to 0 if they are not present.

[0138] If the winding temperature is too low, too much hard tissue will be generated, which will impair the winding performance. Therefore, the preferred lower limit of the temperature is above 350°C.

[0139] Next, after finishing rolling and before annealing the hot-rolled plate, leveling rolling can also be performed. Leveling rolling can correct the shape of the steel plate.

[0140] [Softening Annealing]

[0141] Next, the obtained hot-rolled sheet is annealed before cold rolling to suppress surface decarburization and coarsen the carbides in the steel. This process is crucial for sharpening the crystal orientation after final annealing, i.e., obtaining the aforementioned ODF. When coiled under the coiling temperature conditions of this invention, a large amount of hard microstructure is formed. In this state, the cold rolling load is high, so this process is also useful from the viewpoint of reducing the rolling load.

[0142] [Hot-rolled sheet annealing: carried out for more than 48 hours in a temperature range above 400°C and below Ac1 point, with or without temperature holding.]

[0143] If the annealing temperature exceeds the Ac1 point, an austenitic phase will form, which will result in carbides during the cooling process that are detrimental to the formation of {111} grains in the final annealed sheet. Therefore, the annealing temperature is below the Ac1 point. On the other hand, if the temperature is too low, the effect of carbide coarsening will be smaller, so the lower limit is set at 400°C, and more preferably at 450°C or higher.

[0144] Furthermore, the annealing time is 48 hours or more. This limitation is based on the fact that a long annealing time is necessary to fully spheroidize the slightly fine carbides precipitated during coiling. If the annealing time is less than 48 hours, the spheroidization of the carbides cannot proceed sufficiently, making it difficult to obtain {111} grains that are beneficial to local ductility in subsequent rolling annealing processes. A more preferred annealing time is 60 hours or more. Additionally, decarburization from the steel sheet surface is generally promoted at higher temperatures during annealing, but for the reasons mentioned above, the annealing temperature must be below Ac1, thus suppressing the promotion of decarburization caused by high temperatures. Even so, while decarburization proceeds through long annealing, it is suppressed due to the addition of specified amounts of Cu, Sb, and Sn.

[0145] There is no particular upper limit to the annealing time, but from the perspective of productivity and cost, it is preferable to be below 300 hours. It should be noted that, as long as the annealing temperature is above 400°C and below the Ac1 point, the effect remains the same whether or not a constant temperature is maintained. As an example of not maintaining a constant temperature, the temperature can be raised to above 400°C and below the Ac1 point over 48 hours.

[0146] [Cold Rolled]

[0147] Next, the steel sheet is preferably cold-rolled to the desired thickness after pickling. Cold rolling can be performed using either tandem rolling (unidirectional rolling) or reverse rolling, or it can utilize known warm rolling techniques or inter-pass aging techniques. The reasons for limiting the cold rolling conditions are as follows.

[0148] [Total reduction rate: 50% or more]

[0149] The total reduction ratio is also important for obtaining the aforementioned ODF. When the total reduction ratio is low, the recrystallization driving force is low, so sufficient recrystallization cannot occur in the subsequent annealing process, making it impossible to achieve both high strength and high ductility. In addition, {111} grains, which contribute to improving local ductility, are more easily obtained with a high cold rolling ratio; therefore, the total reduction ratio is 50% or more, more preferably 60% or more. There is no specific upper limit, but from the viewpoint of rolling load, it is preferable to be 95% or less.

[0150] [Rolling with a rolling rate of 30% or more in at least one of the first or second passes]

[0151] In this invention, the process is required to obtain the aforementioned ODF after final annealing. If the rolling rate of the first or second pass is low, strain concentrates on the surface layer, making it difficult to form {111} grains that contribute to improved local ductility within the plate thickness. For the formation of {111} grains, a high rolling rate of 30% or more is preferred for the first pass. A rolling rate of 35% or more is more preferred. There is no specific upper limit, but if it is too high, the steel sheet shape will deteriorate, so it is set to less than 50%, and here it is set to less than 49%. As for the number of rolling passes, one or more passes are sufficient.

[0152] [Two-stage annealing process]

[0153] Next, as an annealing process, the resulting steel sheet undergoes heat treatment and, if necessary, plating. First, the cold-rolled steel sheet is heated to above 600°C but below the Ac1 point. The average heating rate is not specifically limited, but if it is too low, grain coarsening may impair strength; therefore, it is set to above 5°C / s. There is no specific upper limit; conditions exceeding 100°C / s can be applied without issue using methods such as induction heating.

[0154] In the heat treatment, the temperature is first raised to a temperature range above 600°C but below point Ac1 (stage 1), held in this temperature range for at least 60 seconds, and then raised to a temperature range above point Ac1 but below 950°C (stage 2). The average cooling rate is then maintained at at least 10°C / s up to 500°C, cooling to a cooling stop temperature range of 50°C to 350°C. Next, the temperature is raised to a temperature range above the cooling stop temperature but above 250°C and below 600°C, and held in this temperature range for at least 10 seconds.

[0155] First, the annealing conditions in the first section will be explained as follows.

[0156] [Annealing temperature: in the temperature range above 600℃ but below Ac1, with or without holding the temperature, annealing time: 60s or more]

[0157] The annealing in this first paragraph plays a crucial role in the manufacturing method shown in this invention, as it is a process that actively promotes recrystallization in a temperature range below the Ac1 point. It is known that steel sheets subjected to large strain during cold rolling form a large number of {111} grains after recrystallization. Therefore, if recrystallization can be performed before subsequent heating to above the Ac1 point to induce a phase transformation, a large number of {111} grains can be formed. If the annealing temperature is below 600°C, the rolled structure does not recrystallize, leaving behind coarse structures elongated along the rolling direction. If coarse, unrecrystallized structures remain, the proportion of large-angle grain boundaries where austenite grains preferentially form in the structure before the austenite phase transformation decreases, resulting in a reduction in the number of austenite grains formed. On the other hand, if the temperature is above the Ac1 point, austenite grains form during annealing, and elemental distribution progresses between ferrite and austenite grains, resulting in an uneven structure and impairing the ductility of the final annealed sheet. Therefore, to fully achieve recrystallization, the temperature is set to 600°C or higher, more preferably 650°C or higher. It should be noted that as long as the annealing temperature in the first stage is above 600°C and below the Ac1 point, the effect obtained remains the same whether or not the temperature is kept constant. As an example of not keeping the temperature constant, the temperature can be raised to above 600°C and below the Ac1 point over a period of 60 seconds or more.

[0158] Additionally, recrystallization will not occur if the annealing time is too short, so it is set to 60 seconds or more. There is no specific upper limit.

[0159] Furthermore, this annealing process also effectively suppresses surface decarburization by adding specified amounts of Cu, Sb, and Sn. Previously, in steel sheets without the specified amounts of Cu, Sb, and Sn, a soft decarburized layer formed on the surface during prolonged annealing before cold rolling, resulting in a microstructure with relatively few carbides. When this microstructure is subsequently processed by cold rolling, most of the strain accumulates near the ferrite grain boundaries. As a result, in the subsequent first-stage annealing, {111} grains are excessively generated from the vicinity of these ferrite grain boundaries. On the other hand, fewer {100} grains recrystallize from the vicinity of coarse carbides. Therefore, by adding specified amounts of Cu, Sb, and Sn, even with annealing before cold rolling, a certain amount of {100} grains can be ensured on the surface during this annealing process.

[0160] Furthermore, it is not a problem to perform the first stage of annealing on a different production line than the subsequent second stage of annealing. For example, in the case of long-term annealing, the first stage of annealing can be performed in batches, followed by the second stage of annealing in continuous annealing.

[0161] Next, the annealing conditions in the second paragraph will be explained.

[0162] [Annealing temperature: above Ac1 and below 950℃]

[0163] If the annealing temperature is below Ac1, a large amount of ferrite phase remains in the annealed microstructure, resulting in insufficient strength. To obtain high strength, it is preferable to have as little ferrite phase as possible; therefore, the annealing temperature is more preferably above Ac3. On the other hand, if the annealing temperature is above 950°C, the microstructure becomes too large, making it difficult to balance high strength and ductility, and also reducing {111} grains; therefore, it is more preferable to have a temperature below 920°C.

[0164] It should be noted that the Ac1 and Ac3 points shown in this specification can be calculated using the following formulas (A) and (B) as described in "Leslie Steel Materials Science" (Maruzen Co., Ltd., translated by Shigeyasu Koda, published May 31, 1985, page 273). In the formulas, [] represents the content (mass%) of each element. Elements not present in the steel plate or whose content is below the analytical detection limit are calculated as 0% by mass.

[0165] Ac1 point (℃)=723-10.7×[Mn]-16.9×[Ni]+29.1×[Si]+16.9×[Cr]+290×[As]+6.38×[W]……(A)

[0166] Ac3 point (℃)=910-203×√[C]-15.2×[Ni]+44.7×[Si]+104×[V]+31.5×[Mo]+13.1×[ W]-30×[Mn]-11×[Cr]-20×[Cu]+700×[P]+400×[Al]+120×[As]+400×[Ti]……(B)

[0167] [Average cooling rate up to at least 500°C: 10°C / s or more]

[0168] In the cooling process after annealing, a high average cooling rate, preferably 10°C / s or higher, is preferred, up to at least 500°C. This is to suppress the amount of newly formed ferrite due to ferrite phase transformation and to suppress the grain growth of existing ferrite, thus avoiding excessive softening of the material. More preferably, it is 15°C / s or higher. On the other hand, no specific upper limit is specified. The reason for specifying an average cooling rate up to at least 500°C is that, within the composition range of the present invention, a low cooling rate near this temperature region easily leads to the formation of structures such as pearlite and bainite, which are detrimental to improving the balance between strength and ductility.

[0169] [Cooling stop temperature: 50℃~350℃]

[0170] The annealing followed by cooling is carried out at 50°C to 350°C. By cooling to below 350°C, a portion of the austenite produced during annealing is transformed into martensite. If the temperature is below 50°C, all the austenite required for improved ductility is transformed into martensite and disappears. Furthermore, if the temperature exceeds 350°C, the martensitic transformation is insufficient, and the proportion of tempered martensite required for high strength decreases. On the other hand, the proportion of fresh martensite, which is useful for high strength but lacks ductility, increases, making it difficult to achieve both high strength and high ductility even with adjustments to subsequent heat treatment conditions. The preferred lower limit is 75°C or higher. The preferred upper limit is 300°C or lower. The cooling method of this invention can employ known methods, including gas cooling, oil cooling, mist cooling, and cooling of low-melting-point liquid metals.

[0171] [Reheating: After heating to a temperature range above the cooling stop temperature but above 250°C and below 600°C, maintain that temperature range for at least 10 seconds.]

[0172] After the cooling process stops, reheating is performed to temper the generated martensite and stabilize the remaining austenite. Reheating can be performed immediately after cooling stops, or it can be performed after maintaining the temperature within a range where carbide precipitation is not significant. Here, the temperature range where carbide precipitation is not significant is preferably 250°C or lower. More preferably, it is 200°C or lower.

[0173] If the reheating temperature is too low, the diffusion of carbon required to obtain the aforementioned structure cannot be sufficiently promoted, and both high strength and high ductility cannot be achieved. Conversely, if the reheating temperature is too high, coarse carbides precipitate, leading to softening. Therefore, it is necessary to heat to a temperature range above the cooling stop temperature but exceeding 250°C and below 600°C, and then hold this temperature for at least 10 seconds. A more preferred range is above the cooling stop temperature but exceeding 280°C and below 400°C. Furthermore, if the holding time is less than 10 seconds, the diffusion of carbon required for austenite stabilization cannot proceed sufficiently, and ductility cannot be adequately improved. More preferably, 30 seconds or more, and even more preferably, 60 seconds or more.

[0174] Furthermore, in the annealing process, hot-dip galvanizing can be performed within the aforementioned temperature range, i.e., above 250°C and below 600°C, to produce hot-dip galvanized steel sheets. Alternatively, alloying treatment can be performed in a temperature range of 470°C to 600°C after hot-dip galvanizing to produce alloyed hot-dip galvanized steel sheets. Additionally, the steel sheets of this invention can be electroplated to produce electroplated steel sheets.

[0175] Temper rolling can also be performed after annealing. The preferred range for elongation is 0.02% to 2.0%.

[0176] It should be noted that manufacturing conditions other than those mentioned above can be carried out using conventional methods.

[0177] The technology disclosed herein enables the manufacture of thin steel sheets with high strength and excellent ductility.

[0178] Example

[0179] The following describes embodiments of the present invention. However, the present invention is not limited to the embodiments described below, and can be implemented with appropriate modifications within the scope of the spirit of the present invention, all of which are included within the technical scope of the present invention.

[0180] The steel billet, having the composition shown in Table 4 and with the remainder consisting of Fe and unavoidable impurities, is melted, rolled into a billet, and then hot-rolled and cooled under the conditions shown in Table 5 before being coiled. It is then pickled and cold-rolled. Next, it undergoes one or two heat treatments. It should be noted that the number of heat treatments is one if the first heating stage involves continuous heating, and two if the first stage involves discontinuous heating. For some, hot-dip galvanizing, alloying hot-dip galvanizing, and electro-galvanizing are performed in subsequent heat treatment processes.

[0181]

[0182]

[0183]

[0184] [Tension test]

[0185] JIS No. 5 test pieces (marking distance 50 mm, parallel section width 25 mm) were cut from the obtained steel plate in a direction perpendicular to the rolling direction, and tensile tests were performed according to JIS Z 2241. Tensile strength of 980 MPa or higher was considered high strength. Furthermore, the elongation value was calculated by measuring the distance between the markings when the test pieces were butted together after tensile fracture, as the change in elongation before and after the tensile test. Here, elongation of 10% or more was considered high elongation.

[0186] [Organizational Observation]

[0187] In addition, the test pieces were embedded in carbon resin with the rolling direction and thickness direction as observation surfaces for microstructure evaluation.

[0188] [Hydrogen embrittlement characteristics test of the bending part]

[0189] In addition, a 50L×100C sample was punched out from 1 / 4 of the plate width with a 15% gap, and then bent at 50° to form a V shape, resulting in a bending R / t = 5, thus preparing a test piece. Next, the opening angle of the bent portion of the test piece was varied, and the surface of the test piece was constrained under a residual stress of up to 1.5×TS. The constrained test piece was then placed in a hydrogen immersion environment of hydrochloric acid at a liquid temperature of 25°C and pH 0.9 for 24 hours, and the presence of cracks in the central part of the bent portion was visually observed. Test pieces without visually confirmed cracks were judged to have excellent hydrogen embrittlement resistance in the bent portion.

Claims

1. A high-strength steel plate, comprising the following components: It contains, by mass%, C: 0.090%–0.300%, Si: 0.40%–2.50%, Mn: 1.8%–4.0%, P: less than 0.100%, S: less than 0.0200%, Al: less than 0.200%, N: less than 0.0200%, and O: less than 0.0100%. It contains at least two of the following: Cu: 0.005%–0.5%, Sn: 0.005%–0.5%, and Sb: 0.001%–0.07%, totaling 0.01%–1.0%. It contains at least one of the following: Ti: less than 0.200%, Nb: less than 0.200%, and V: less than 0.200%, with a content of 0.01% or more. The remainder consists of Fe and unavoidable impurities; The microstructure of the steel at the 1 / 4 position of the plate thickness satisfies: The total area of ​​ferrite and bainite ranges from 0% to 20%. The retained austenite content is 4%–20% by volume. The tempered martensite content is over 70% by area ratio. Fresh martensite has an area ratio of 0% to 20%, and The microstructure of the outermost layer of steel satisfies: The total area of ​​ferrite and bainite ranges from 0% to 20%. The retained austenite content is 4%–20% by volume. The tempered martensite content is over 70% by area ratio. Fresh martensite ranges from 0% to 20% in terms of area. The thickness section of the rolled plate was evaluated using the SEM-EBSD method. The texture of the tempered martensite was determined in the orientation distribution function ODF shown by the Bunge method in the Eulerian space of 0°≤φ1≤90°, φ2=45°, 0°≤Φ≤90°. For the 1 / 4 position of the plate thickness, the maximum ODF of the {111} grain, expressed as the maximum strength at φ1=5×n°, φ2=45°, Φ=55° (n=0, 1, 2, …, 18), is greater than 2.

0. For the outermost layer, the maximum ODF of the {100} grains, expressed as the maximum intensity at φ1=5×n°, φ2=45°, Φ=0° (n=0, 1, 2, …, 18), is greater than 1.

0. The tensile strength of the high-strength steel plate is above 980 MPa.

2. The high-strength steel plate according to claim 1, wherein, The composition, by mass%, further contains at least one element selected from the following: Ta: less than 0.10%, W: less than 0.10%, B: less than 0.0100%, Cr: less than 1.00%, Mo: less than 1.00%, Co: less than 0.010%, Ni: less than 1.00%, Ca: less than 0.0100%, Mg: less than 0.0100%, REM: less than 0.0100%, Zr: less than 0.100%, Te: less than 0.100%, Hf: less than 0.10%, and Bi: less than 0.200%.

3. The high-strength steel plate according to claim 2, wherein, The composition of the ingredients, by mass%, also contains Ti: 0.010%–0.200%, Nb: 0.005%–0.200%, and B: 0.0005%–0.0100%.

4. The high-strength steel plate according to any one of claims 1 to 3, wherein, An electroplated zinc layer is formed on the surface.

5. The high-strength steel plate according to any one of claims 1 to 3, wherein, A hot-dip galvanized layer is formed on the surface.

6. The high-strength steel plate according to any one of claims 1 to 3, wherein, An alloyed hot-dip galvanized layer is formed on the surface.

7. A method for manufacturing a high-strength steel plate, as described in any one of claims 1 to 3. After heating the steel billet with the aforementioned composition to a temperature range of 1080°C to 1300°C Finish rolling is performed at the finishing temperature within the temperature range of 850℃ to 1000℃. After finishing rolling, the temperature is cooled to below 720°C within 2 seconds at an average cooling rate of over 70°C per second. After winding in a temperature range below Ms+100℃ Prior to cold rolling, the material undergoes softening annealing for at least 48 hours in a temperature range above 400°C and below Ac1, with or without temperature holding. Then, cold rolling is performed, wherein at least one of the first or second passes has a rolling rate of 30% or more, and the total reduction rate is 50% or more. This process yields a cold-rolled sheet. The cold-rolled sheet is annealed for at least 60 seconds in a temperature range above 600°C and below the Ac1 point, with or without holding the temperature. Then, heat to above point Ac1 and below 950°C. Furthermore, the average cooling rate up to at least 500°C is 10°C / s or higher. Cooling to the cooling stop temperature range of 50℃ to 350℃ Next, heat to a temperature range above the cooling stop temperature but above 250°C and below 600°C, and maintain that temperature range for at least 10 seconds.

8. The method for manufacturing high-strength steel plate according to claim 7, wherein, Further implement electroplating zinc.

9. The method for manufacturing high-strength steel plate according to claim 7, wherein, Further implement hot-dip galvanizing.

10. The method for manufacturing high-strength steel plate according to claim 7, wherein, Further hot-dip galvanizing is carried out, followed by alloying treatment of galvanizing in a temperature range of 470℃~600℃.

Images