JP7666761B1 - High strength steel plate and method for manufacturing same - Google Patents

Abstract

The object of the present invention is to provide a high-strength steel plate having a tensile strength of 980 MPa or more and excellent ductility, fatigue strength and punchability, and a method for manufacturing the same.
1. A high-strength steel plate having a predetermined chemical composition, characterized in that a bainite phase consisting of upper bainite and granular bainite in an area ratio of 85% or more is a main phase, the area ratio of a hard second phase containing fresh martensite and/or a retained austenite phase is less than 15%, the area ratio of granular bainite relative to the area ratio of upper bainite in the bainite phase is 5% or more and 40% or less, the difference |Hv1-Hv2| between a hardness Hv1 at 1/2 the plate thickness position and a hardness Hv2 at 1/4 the plate thickness position is 30% or less for 0.3TS, and the upper bainite phase has an average crystal grain size of 7 μm or less.

Description

The present invention relates to a high-strength steel sheet (particularly a hot-rolled steel sheet) suitable for use as an automotive component, which has improved fatigue resistance and punchability in addition to strength and ductility, and a method for manufacturing the same.

In recent years, from the viewpoint of protecting the global environment, there has been a demand for a reduction in _CO2 emissions on a global framework. In particular, there is a strong demand for improved fuel efficiency in automobiles, and there is a trend toward reducing the weight of automobile bodies. Increasing the strength of steel plates, which are the raw material for automobile parts, and reducing their thickness is one effective means of reducing the weight of automobile bodies without decreasing their strength. In particular, steel plates with a tensile strength of 980 MPa or more are expected to be a material that dramatically improves automobile fuel efficiency through thinner thickness.

On the other hand, increasing the tensile strength of steel sheet reduces its ductility, which deteriorates its press formability. Automobile parts, particularly suspension parts and other undercarriage parts, need to be made into complex shapes to ensure rigidity, so materials for these parts require high press formability, and in particular excellent ductility.

Furthermore, in order to ensure durability, which decreases as automotive parts become thinner, it is necessary to improve the fatigue strength of steel plates. Automotive parts, particularly suspension parts and other undercarriage parts, are subjected to repeated loads from tires, so if fatigue strength is low, the durability of the parts may fall short of what was expected in the design as the driving distance increases. However, generally speaking, increasing the strength of steel plates does not necessarily increase fatigue strength, and in addition, increasing the strength of materials makes it easier for defects with unevenness to occur on the punched end surface, which then initiates cracks, leading to a deterioration in fatigue resistance.

Thus, in steel sheets with a tensile strength of 980 MPa or more, excellent fatigue strength and punchability are required in addition to excellent ductility. Various studies have been conducted to improve the fatigue strength and punchability while increasing the tensile strength of steel sheets (Patent Documents 1 to 3).

Patent Document 1 discloses technology relating to high-strength hot-rolled steel sheet with excellent formability and fatigue resistance by controlling the manufacturing conditions of hot rolling, making the main phase ferrite, and controlling the shape and dispersion form of inclusions.

Patent Document 2 discloses technology relating to high-strength hot-rolled steel sheet that has excellent punching fatigue resistance and workability by controlling the shape and hardness of martensite in the center of the sheet thickness.

Patent Document 3 discloses technology relating to a high-strength hot-rolled steel sheet with excellent hole expandability and punchability, which is achieved by controlling the manufacturing conditions of the hot rolling, making the main phase bainite, and controlling the shape and dispersion form of the hard second phase as well as the amount of precipitates.

JP 2014-31560 A JP 2015-214718 A Republished Publication No. 2017-017933

However, the conventional techniques described in Patent Documents 1 to 3 had the following problems.

The technologies described in Patent Documents 1 and 2 do not achieve a tensile strength of 980 MPa or more.

The technology described in Patent Document 3 provides a hot-rolled steel sheet having a tensile strength of 980 MPa or more, excellent fatigue strength, and excellent punchability. The technology described in Patent Document 3 provides a hot-rolled steel sheet having a tensile strength of 980 MPa or more, excellent punchability, and excellent hole expandability. However, there is no mention of the fatigue strength that is the subject of this application, and it is not necessarily the case that good fatigue strength is obtained.

As such, conventional technology does not have established technology for producing high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent ductility, fatigue strength, and punchability.

Therefore, the present invention aims to provide a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent ductility, fatigue strength and punchability, and a manufacturing method thereof.

In the present invention, excellent ductility means that the steel has a uniform elongation of 6.0% or more. Excellent fatigue strength means that the ratio of plane bending fatigue strength at 2×10 ⁶ times to the tensile strength (fatigue limit ratio) is 0.50 or more in a plane bending test with a stress ratio of -1. Excellent punchability means that when punched with a 10 mmφ punch at three or more clearances including 10% and 20% within a clearance range of 10 to 20%, there are no cracks, chips, brittle fracture surfaces, or secondary shear surfaces on the end surfaces of each punched hole.

In order to achieve the above object, the inventors have conducted intensive research to improve the ductility, fatigue strength, and punchability of hot-rolled steel sheets while maintaining a tensile strength of 980 MPa or more. As a result, the main phase is bainite, and the area ratio of soft granular bainite and upper bainite generated near the Bs point is controlled within a certain range, and the hard second phase, fresh martensite and retained austenite phase, is set to less than 15%. In addition, segregation is reduced by setting the difference |Hv1-Hv2| between the hardness Hv1 at 1/2 the sheet thickness position and the hardness Hv2 at 1/4 the sheet thickness position to 30% or less for 0.3TS (MPa). Furthermore, the average crystal grain size of the upper bainite is set to 7 μm or less. It has been found that these make it possible to obtain a steel sheet having excellent ductility, fatigue strength, and punchability while maintaining a tensile strength of 980 MPa or more.

Here, the upper bainite phase is an aggregate of lath-shaped bainitic ferrite. Specifically, it means a structure having Fe-based carbides and/or retained austenite phase between lath-shaped bainitic ferrite.

Lathyrus bainitic ferrite differs from lamellar (layered) ferrite and polygonal ferrite in pearlite in that it has a lathyrus shape and a relatively high dislocation density internally, making the two distinct from each other using a scanning electron microscope (SEM) or transmission electron microscope (TEM).

In addition, when there is a residual austenite phase between the laths, only the lath-shaped bainitic ferrite portion is considered as upper bainite and is distinguished from the residual austenite phase. Fresh martensite is martensite that does not contain Fe-based carbides.

Granular bainite is composed of granular bainitic ferrite and means a structure having Fe-based carbides and/or residual austenite phase between the granular bainitic ferrite. The aspect ratio is 2 or less, and the grains have grain boundaries that do not have the orientation relationship of <axial direction>-rotation angle of <15 16 18>-60°, <6 15 15>-56°, <6 7 7>-50°, <1 0 1>-60°, <1 2 2>-17°. The crystal orientation is determined using the Electron Backscatter Diffraction Patterns (EBSD) method. Fresh martensite and/or residual austenite phase have low clarity (IQ value). Prior austenite grains have a large aspect ratio and large crystal grains, so in this study, crystal grains with an aspect ratio of 2 or more and a size of 10 μm or more were considered to be prior austenite grains. Therefore, upper bainite grains were excluded from crystal grains with an orientation difference of 15° or more, and of the remaining crystal grains, crystal grains with a high IQ value and an aspect ratio of 2 or less for 10 μm or less were determined to be granular bainite.

In addition, the fresh martensite and/or retained austenite phase has brighter contrast in the SEM image than the upper bainite phase, the lower bainite and/or tempered martensite phase, and the polygonal ferrite phase. Therefore, the fresh martensite phase and/or retained austenite phase can be distinguished from these structures using SEM.

The fresh martensite and retained austenite phases have similar contrast in SEM, but can be distinguished from each other using electron beam reflection diffraction techniques.

The present invention was made based on the above findings and further studies, and is summarized as follows.
[1] A steel sheet having a composition, by mass%, of C: 0.04-0.18%, Si: 0.1-2.5%, Mn: 1.0-3.0%, P: 0.1% or less, S: 0.01% or less, Al: 0.010-1.0%, and N: 0.01% or less, with the balance being Fe and unavoidable impurities. The microstructure has a bainite phase consisting of upper bainite and granular bainite at an area ratio of 85% or more as the main phase, and fresh martensitic a hard second phase including a bainite phase and/or a retained austenite phase having an area ratio of less than 15%, an area ratio of granular bainite relative to the area ratio of upper bainite in the bainite phase being 5% or more and 40% or less, a difference |Hv1-Hv2| between a hardness Hv1 at a position 1/2 of the plate thickness and a hardness Hv2 at a position 1/4 of the plate thickness being 30% or less for 0.3TS, and an average crystal grain size of the upper bainite phase being 7 μm or less.
[2] The high-strength steel plate according to [1], further comprising, in addition to the above-mentioned composition, one or more of groups a to e in mass %, Group a: one or more selected from V: 0.005-0.5%, Ti: 0.005-0.2%, and Nb: 0.005-0.1%, Group b: one or more selected from Cu: 0.005-0.5%, Ni: 0.005-0.5%, Cr: 0.005-1.0%, and Mo: 0.005-0.5%, Group c: B: 0.0002-0.005%, group d: Sb: 0.001-0.1%, and Sn: 0.001-0.1%, one or two selected from group e: Ca: 0.0001-0.005%, Mg: 0.0001-0.005%, and REM: 0.0001-0.005%, one or two or more selected from group e.
[3] A method for producing a high-strength steel plate according to [1] or [2], comprising the steps of: casting a molten steel material having the above-mentioned composition while rotating it in a horizontal plane relative to a mold at a rotation speed of 10 cm/s or more by an induction electromagnetic stirrer; heating the obtained steel material to 1150°C or more; rough rolling the heated steel material to obtain a steel plate; finish rolling the steel plate under a finish rolling end temperature condition of (RC-50)°C or more and (RC+100)°C or less; and cooling the steel plate after the finish rolling from the end of the finish rolling. A method for producing a high strength steel sheet, comprising cooling under the conditions of: time to start: 2.0 s or less, average cooling rate on the surface to Bs: 30°C/s or more, residence time at a temperature of (Bs-100)°C or more and Bs°C or less: 3.0 s or more and 10.0 s or less, cooling stop temperature: (Bs-250)°C or more and (Bs-100)°C or less, coiling the cooled steel sheet under the conditions of a coiling temperature: (Bs-250)°C or more and (Bs-100)°C or less, and cooling to (Bs-400)°C or less at an average cooling rate of 1°C/s or less. Note that RC and Bs are defined by the following formulas (1) and (2), respectively.
RC (℃) = 750 + 120 x C + 100 x N + 10 x Mn + 250 x Ti + 5000 x B + 10 x Cr + 50 x Mo + 750 x Nb + 150 x V... (1)
Bs (°C) = 830-270 x C-90 x Mn-70 x Cr-37 x Ni-83 x Mo-20 x Cu... (2)
Here, each element symbol in the above formulas (1) and (2) represents the content (mass %) of each element, and when no element is contained, it is set to 0.

According to the present invention, it is possible to provide a high-strength steel plate having a tensile strength of 980 MPa or more and excellent ductility, fatigue strength, and punchability, and a manufacturing method thereof.

When the high-strength steel plate of the present invention is applied to automobile undercarriage parts such as suspensions, structural parts, skeletal parts, and truck frame parts, it ensures safety and reduces the weight of the automobile body, thereby providing significant industrial benefits.

FIG. 1 is a schematic diagram showing the shape of a test piece for a plane bending fatigue test in the present invention.

The following describes an embodiment of the present invention. Note that the following description shows an example of a preferred embodiment of the present invention, and the present invention is not limited to the following embodiment.

The steel plate has the following chemical composition. In the following description, "%", which is the unit of element content in the chemical composition, means "mass %" unless otherwise specified.

<C: 0.04-0.18%>
C is an element effective in promoting the formation of bainite by improving hardenability and improving strength. It also suppresses the formation of granular bainite by lowering the Bs point and transforming bainite at a lower temperature. If the C content is less than 0.04%, such effects are not sufficiently obtained, and a tensile strength of 980 MPa or more cannot be obtained. For this reason, the C content is 0.04% or more, preferably 0.05% or more, and more preferably 0.06% or more. On the other hand, if the C content exceeds 0.18%, a hard second phase including fresh martensite and/or retained austenite phase is excessively formed, and sufficient fatigue strength cannot be obtained. For this reason, the C content is 0.18% or less, preferably 0.17% or less, and more preferably 0.15% or less.

<Si: 0.1 to 2.5%>
Si strengthens the steel by solid solution and contributes to improving the strength of the steel. Therefore, the Si content is 0.1% or more, preferably 0.3% or more, and more preferably 0.5% or more. On the other hand, Si is an element that promotes the formation of ferrite, and if the Si content exceeds 2.5%, ferrite is formed and the fatigue strength decreases. Therefore, the Si content is 2.5% or less, preferably 2.3% or less, and more preferably 2.0% or less.

<Mn: 1.0 to 3.0%>
Mn is an element that stabilizes austenite and is effective in suppressing the formation of ferrite and improving strength. In addition, it lowers the Bs point and transforms bainite at a lower temperature, thereby suppressing the formation of granular bainite. If the Mn content is less than 1.0%, such effects cannot be sufficiently obtained, and ferrite and granular bainite are formed, and a tensile strength of 980 MPa or more and excellent fatigue strength cannot be obtained. For this reason, the Mn content is 1.0% or more, preferably 1.2% or more, and more preferably 1.5% or more. On the other hand, if the Mn content exceeds 3.0%, the hard second phase including fresh martensite and/or retained austenite phase increases, and sufficient fatigue strength cannot be obtained. For this reason, the Mn content is 3.0% or less, preferably 2.8% or less, and more preferably 2.5% or less.

<P: 0.1% or less>
Since P deteriorates weldability, it is desirable to reduce the amount as much as possible. In the present invention, a P content of up to 0.1% is permissible. Therefore, the P content is set to 0.1% or less. Although there is no particular lower limit, a P content of less than 0.003% leads to an increase in refining costs, so the P content is preferably 0.003% or more, and more preferably 0.005% or more.

<S: 0.01% or less>
Since S deteriorates weldability, it is preferable to reduce the amount as much as possible, but in the present invention, an S content of up to 0.01% is acceptable. Therefore, the S content is set to 0.01% or less. There is no particular lower limit, but since an S content of less than 0.0001% leads to a decrease in production efficiency, it is preferable that the S content is 0.0001% or more, and more preferably 0.0005% or more.

<Al: 0.010-1.0%>
Al acts as a deoxidizer and is an effective element for improving the cleanliness of steel. If the amount of Al is too small, the effect is not necessarily sufficient. Therefore, the Al content is 0.010% or more, preferably 0.015% or more, and more preferably 0.020% or more. On the other hand, Al is an element that promotes the formation of ferrite, and if the Al content exceeds 1.0%, ferrite is generated and the fatigue strength is reduced. Therefore, the Al content is 1.0% or less, preferably 0.8% or less, and more preferably 0.5% or less.

<N: 0.01% or less>
N combines with elements that form nitrides to precipitate as nitrides, contributing to the refinement of crystal grains. However, N is likely to combine with Ti at high temperatures to form coarse nitrides, and excessive N content reduces fatigue strength. For this reason, the N content is 0.01% or less, preferably 0.008% or less, and more preferably 0.006% or less. There is no particular lower limit, but an N content of less than 0.001% leads to a decrease in production efficiency, so 0.001% or more is preferable.

The balance is Fe and unavoidable impurities.

The above components are the basic composition of the high-strength steel plate of the present invention. If necessary, the following elements may be further contained.

V, Ti, and Nb are effective elements for forming carbides and improving strength through precipitation strengthening. For this reason, when V, Ti, and Nb are contained, the respective contents are preferably V: 0.005-0.5%, Ti: 0.005-0.2%, and Nb: 0.005-0.1%. If the V, Ti, and Nb contents exceed the respective upper limits, the carbides may become coarse and the punchability may deteriorate. For this reason, the V content is preferably 0.05% or more, more preferably 0.1% or more, and more preferably 0.3% or less. The Ti content is more preferably 0.01% or more, and more preferably 0.1% or less. The Nb content is more preferably 0.01% or more, and more preferably 0.08% or less.

Cr, Ni, Cu, and Mo are elements that stabilize austenite and are effective in suppressing the formation of ferrite and improving strength. In addition, by lowering the Bs point and transforming bainite at a lower temperature, the formation of granular bainite is suppressed. For this reason, when Cr, Ni, Cu, and Mo are contained, the respective contents are preferably Cr: 0.005-1.0%, Ni: 0.005-0.5%, Cu: 0.005-0.5%, and Mo: 0.005-0.5%. If the contents of Cr, Ni, Cu, and Mo exceed the respective upper limits, a hard second phase containing fresh martensite and/or a retained austenite phase may be excessively formed, and the steel structure of the present invention may not be obtained. The Cr content is preferably 0.01% or more, more preferably 0.3% or more, and more preferably 0.8% or less. The Ni content is preferably 0.01% or more, more preferably 0.05% or more, and more preferably 0.3% or less. The Cu content is preferably 0.01% or more, more preferably 0.05% or more, and more preferably 0.3% or less. The Mo content is preferably 0.01% or more, more preferably 0.05% or more, and more preferably 0.3% or less.

When B is contained, it is preferable to set it to 0.0002 to 0.005%. B is an effective element for segregating at austenite grain boundaries and suppressing the formation of ferrite, thereby promoting the formation of upper bainite and granular bainite and further improving the strength of the steel plate. In order to realize these effects, it is preferable to set the B content to 0.0002% or more. For this reason, the B content is 0.0002% or more, preferably 0.0005% or more, and more preferably 0.0007% or more. On the other hand, when the B content exceeds 0.005%, the above effects are saturated. Therefore, the B content is 0.005% or less, preferably 0.004% or less, and more preferably 0.003% or less.

Sb is an element that is effective in preventing the de-elementation from the surface of the steel material when the steel material is heated, and thus preventing a decrease in the strength of the steel. For this reason, when Sb is contained, the content is preferably set to 0.001 to 0.1%. If the Sb content exceeds the above upper limit, it may lead to embrittlement of the steel plate. The Sb content is more preferably 0.005% or more, and more preferably 0.05% or less.

Sn is an element that is effective in suppressing the formation of pearlite and preventing a decrease in the strength of steel. In order to obtain this effect, when Sn is contained, the content is preferably set to 0.001 to 0.1%. If the Sn content exceeds the above upper limit, it may lead to embrittlement of the steel plate. The Sn content is more preferably 0.005% or more, and more preferably 0.05% or less.

Ca, Mg, and REM are elements that are effective in improving punchability by controlling the shape of inclusions. In order to obtain such effects, when Ca, Mg, and REM are contained, the respective contents are preferably set to Ca: 0.0001-0.005%, Mg: 0.0001-0.005%, and REM: 0.0001-0.005%. On the other hand, if the Ca and REM contents exceed the above upper limits, the amount of inclusions increases and fatigue strength may deteriorate. The Ca content is more preferably 0.0005% or more and more preferably 0.003% or less. The Mg content is more preferably 0.0005% or more and more preferably 0.003% or less. The REM content is more preferably 0.0005% or more and more preferably 0.003% or less. Note that REM (rare earth elements) is a collective term for Sc, Y, and 15 elements ranging from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content here refers to the total content of these elements. Note that for each component of groups a to e (V, Ti, Nb, Cr, Ni, Cu, Mo, B, Sb, Sn, Ca, Mg, REM), when the content is less than the lower limit value mentioned above, the component is considered to be included as an inevitable impurity.

Next, we will explain the microstructure of the high-strength steel plate of the present invention.

The high-strength steel plate of the present invention has the following microstructure. The main phase is a bainite phase consisting of upper bainite and granular bainite with an area ratio of 85% or more. The area ratio of a hard second phase including fresh martensite and/or retained austenite phase is less than 15%. In the bainite phase, the area ratio of granular bainite relative to the area ratio of upper bainite is 5% or more and 40% or less, and the average crystal grain size of the upper bainite phase is 7 μm or less. The difference |Hv1-Hv2| between the hardness Hv1 at 1/2 the plate thickness position and the hardness Hv2 at 1/4 the plate thickness position is 30% or less for 0.3TS.

<Main phase: 85% or more bainite phase>
The microstructure of the high-strength steel plate of the present invention contains a bainite phase consisting of upper bainite and granular bainite as a main phase. If the bainite phase is less than 85% by area, excellent punchability cannot be obtained. Therefore, the bainite phase is set to 85% or more, preferably 90% or more, and more preferably 95% or more.

<Hard second phase: Fresh martensite and/or retained austenite phase less than 15%>
If the hard second phase is controlled to be fresh martensite and/or retained austenite and the area ratio is less than 15%, no macroscopic stress concentration occurs at the phase interface during fatigue testing, and no fatigue cracks will be initiated, resulting in excellent fatigue strength. For this reason, the area ratio of the hard second phase containing fresh martensite and/or retained austenite is set to less than 15%, preferably 10% or less, and more preferably 5% or less.

In addition, the remaining structure other than upper bainite, granular bainite, fresh martensite and/or retained austenite phase is acceptable if it is 3% or less. In addition, examples of the remaining structure include lower bainite, ferrite, tempered martensite, pearlite, etc.

<Area ratio of granular bainite to area ratio of upper bainite in the bainite phase is 5% to 40%>
Granular bainite has better ductility than upper bainite. When the area ratio of granular bainite to the area ratio of upper bainite in the bainite phase is less than 5%, excellent ductility cannot be obtained. Therefore, the area ratio of granular bainite to the area ratio of upper bainite in the bainite phase is 5% or more, preferably 8% or more, and more preferably 10% or more. On the other hand, when the area ratio of granular bainite to the area ratio of upper bainite exceeds 40%, a significant deterioration in fatigue strength is caused. Therefore, the area ratio of granular bainite to the area ratio of upper bainite is 40% or less, preferably 35% or less, and more preferably 30% or less.

<Average grain size of upper bainite phase is 7 μm or less>
The grain boundaries have the effect of making it difficult for the slip deformation to propagate to adjacent grains, and as a result, the fatigue strength can be improved. Therefore, the average grain size of the upper bainite phase is 7 μm or less, and preferably 6 μm or less. Although there is no particular lower limit, if the average grain size is too small, it may lead to a decrease in ductility, so the average grain size is preferably 1 μm or more, and more preferably 2 μm or more.

<The difference between the hardness Hv1 at 1/2 the plate thickness position and the hardness Hv2 at 1/4 the plate thickness position |Hv1-Hv2| is 30% or less for 0.3TS>
When the segregation of components such as Mn becomes excessive at the 1/2 position of the sheet thickness, a hard second phase containing fresh martensite and/or a retained austenite phase is locally generated, and a difference in hardness occurs with the parent phase structure, which makes it easy for cracks and roughness to occur during punching. When the difference |Hv1-Hv2| between the hardness Hv1 at the 1/2 position of the sheet thickness and the hardness Hv2 at the 1/4 position of the sheet thickness exceeds 30% with respect to 0.3TS, the punchability is significantly deteriorated. Therefore, the difference |Hv1-Hv2| between the hardness Hv1 at the 1/2 position of the sheet thickness and the hardness Hv2 at the 1/4 position of the sheet thickness is 30% or less with respect to 0.3TS, preferably 25% or less, and more preferably 20% or less. The lower limit does not need to be particularly limited, and may be 0%. TS is tensile strength. The measurements of the hardness Hv1, Hv2, and TS at this time are performed by the method shown in the examples.

The high-strength steel plate of the present invention has a tensile strength of 980 MPa or more, and also has excellent ductility, excellent fatigue strength, and excellent punchability. Therefore, the high-strength steel plate of the present invention has high tensile strength, and even when thinned, it ensures safety and can be applied to components of trucks and passenger cars. In addition, in the present invention, the area ratios of each of the above-mentioned structures and the mechanical properties are values measured by the method described in the examples.

Next, a manufacturing method of a high-strength steel plate according to one embodiment of the present invention will be described. In the following description, the temperature indication "℃" refers to the surface temperature of the object (steel material or steel plate) unless otherwise specified.

The high-strength steel plate of the present invention can be produced by sequentially carrying out the following processes (1) to (6). Each process will be described below.
(1) Melting and casting (2) Heating (3) Hot rolling (4) Cooling (first cooling)
(5) Winding (6) Cooling (second cooling)
(1) Melting and Casting The method of melting the molten steel is not particularly limited. For example, the molten steel having the above-mentioned composition can be melted by a known method such as a converter to obtain the molten steel. Scrap may be used as the raw material. The obtained molten steel is made into a steel material by a continuous casting method. Induction electromagnetic stirring has the effect of reducing the segregation of components that occurs during solidification by dividing the columnar crystals growing from the mold and generating equiaxed crystals from the 1/4 to 1/2 plate thickness positions, and promotes the formation of granular bainite. If the rotation speed during electromagnetic stirring is less than 10 cm/s, the columnar crystals are not sufficiently divided and such an effect cannot be obtained. For this reason, the rotation speed is 10 cm/s or more in the horizontal plane, preferably 15 cm/s or more, and more preferably 20 cm/s or more.

On the other hand, the upper limit of the swirling speed during electromagnetic stirring is not particularly limited, but from the viewpoint of production, it is preferably 50 cm/s or less, more preferably 45 cm/s or less, and still more preferably 40 cm/s or less. After being produced by the continuous casting method, the steel material may be directly subjected to the subsequent heating step, or the steel material that has been cooled to become a hot or cold piece may be subjected to the heating step.
Furthermore, the chemical composition of the finally obtained high strength steel plate is the same as that of the raw steel material used.

(2) Heating First, the steel material is heated to a heating temperature of 1150°C or more. In the steel material after being cooled to a low temperature, most of the precipitate-forming elements are present in a non-uniform manner as coarse precipitates. If the additive elements are present as coarse and non-uniform precipitates, the effect of the additive elements is not fully obtained, and the desired steel structure is not obtained. Therefore, it is necessary to heat the steel material prior to hot rolling to dissolve the coarse precipitates. For this reason, the heating temperature of the steel material is 1150°C or more, preferably 1180°C or more, and more preferably 1200°C or more. On the other hand, if the heating temperature of the steel material becomes too high, it will cause the occurrence of slab defects and a decrease in yield due to scale-off. For this reason, the heating temperature of the steel material is preferably 1350°C or less, more preferably 1300°C or less, and even more preferably 1280°C or less. In terms of uniformizing the temperature of the steel material during heating, it is preferable to heat the steel material to the heating temperature and then hold it at the heating temperature. The time for holding at the heating temperature (holding time) is not particularly limited, but from the viewpoint of increasing the temperature uniformity of the steel material, it is preferable to set it to 1800 seconds or more. On the other hand, if the holding time exceeds 10000 seconds, the amount of scale generation increases. As a result, scale bite or the like is likely to occur in the subsequent hot rolling, leading to a decrease in yield due to poor surface defects. Therefore, the holding time is preferably 10000 seconds or less, and more preferably 8000 seconds or less. Note that this heating step also includes a case where the steel material before hot rolling is directly subjected to hot rolling (direct rolling) while still at a high temperature (i.e., while maintaining the temperature in the above heating temperature range) after casting.

(3) Hot rolling Next, the heated steel material (or the steel material still at a high temperature after casting) is subjected to hot rolling consisting of rough rolling and finish rolling. The conditions of the rough rolling are not particularly limited as long as the desired sheet bar dimensions can be secured.

First, the steel material is roughly rolled to obtain a roughly rolled plate. Before the resulting roughly rolled plate is subjected to finish rolling, it may be subjected to descaling (high-pressure water descaling) in which high-pressure water is sprayed on the entry side of the finish rolling mill.

Next, in the present invention, when the temperature RC is defined by the following formula (1), the finish rolling end temperature is set to be (RC-50)°C or more and (RC+100)°C or less. RC is the lower limit temperature of austenite recrystallization estimated from the component composition. If the finish rolling end temperature is less than (RC-50)°C, strain relief by recrystallization is difficult to occur, and the austenite continues to accumulate strain as it transitions to cooling, excessively promoting the formation of granular bainite. For this reason, the finish rolling end temperature is (RC-50)°C or more, preferably (RC-30)°C or more, and more preferably RC°C or more.

On the other hand, if the finish rolling end temperature is higher than (RC+100)°C, the austenite grains become coarse and the average grain size of the upper bainite becomes large, so sufficient fatigue strength cannot be obtained. Therefore, the finish rolling end temperature is (RC+100)°C or less, preferably (RC+80)°C or less, and more preferably (RC+50)°C or less. RC is defined by the following formula (1).
RC (℃) = 750 + 120 x C + 100 x N + 10 x Mn + 250 x Ti + 5000 x B + 10 x Cr + 50 x Mo + 750 x Nb + 150 x V... (1)
Here, each element symbol in the above formula (1) represents the content (mass%) of each element, and when the element is not contained, it is set to 0.

(4) Cooling (first cooling)
Next, the obtained hot-rolled steel sheet is cooled (first cooling). In this case, the time from the end of hot rolling to the start of cooling (cooling start time) is set to within 2.0 seconds after the end of finish rolling. If the cooling start time exceeds 2.0 seconds, the austenite grains become coarse and the average grain size of the upper bainite becomes large, so that sufficient tensile strength and fatigue strength cannot be obtained. Therefore, the cooling start time is set to within 2.0 seconds, preferably within 1.5 seconds, and more preferably within 1.0 seconds.

In addition, the average cooling rate on the surface from the finish rolling end temperature to Bs is 30°C/s or more. If the average cooling rate from the finish rolling end temperature to Bs is too slow, ferrite is generated excessively, and sufficient tensile strength and fatigue strength cannot be obtained. For this reason, the average cooling rate is 30°C/s or more, preferably 40°C/s or more, and more preferably 50°C/s or more. On the other hand, although there is no particular upper limit, if it is too fast, it becomes difficult to manage the cooling stop temperature. For this reason, the average cooling rate is preferably 500°C/s or less, more preferably 300°C/s or less, and even more preferably 150°C/s or less.

Furthermore, during cooling, the residence time at a temperature of (Bs-100)°C or higher and Bs°C or lower is set to 3.0 s or higher and 10.0 s or lower. If the residence time at a temperature of (Bs-100)°C or higher and Bs°C or lower is less than 3.0 s, sufficient granular bainite is not obtained. For this reason, the residence time at a temperature of (Bs-100)°C or higher and Bs°C or lower is 3.0 s or higher, preferably 4.0 s or higher, and more preferably 5.0 s or higher. If the residence time at a temperature of (Bs-100)°C or higher and Bs°C or lower exceeds 10.0 s, excessive granular bainite is generated and sufficient fatigue strength is not obtained. For this reason, the residence time at a temperature of (Bs-100)°C or higher and Bs°C or lower is set to 10.0 s or lower, preferably 9.0 s or lower, and more preferably 8.0 s or lower.

In addition, the cooling method is not particularly limited as long as forced cooling is performed so as to obtain the above average cooling rate. The cooling stop temperature is (Bs-250)°C or more and (Bs-100)°C or less. If the cooling stop temperature is less than (Bs-250)°C, the microstructure becomes lower bainite. The lower bainite is a high-strength structure, but does not have excellent ductility. For this reason, the cooling stop temperature is (Bs-250)°C or more, preferably (Bs-220)°C or more, and more preferably (Bs-200)°C or more. On the other hand, if the cooling stop temperature is higher than (Bs-100)°C, granular bainite is excessively generated, and therefore excellent fatigue strength cannot be obtained. For this reason, the cooling stop temperature is (Bs-100)°C or less, preferably (Bs-120)°C or less, and more preferably (Bs-150)°C or less. Note that Bs is defined by the following formula (2).
Bs (℃) = 830-270 x C-90 x Mn-70 x Cr-37 x Ni-83 x Mo-20 x Cu... (2)
Here, each element symbol in the above formula (2) represents the content (mass %) of each element, and an element that is not contained is set to 0.

(5) Coiling Next, the cooled hot-rolled steel sheet is coiled under the condition of a coiling temperature of (Bs-250)°C or more and (Bs-100)°C or less. If the coiling temperature is less than (Bs-250)°C, lower bainite and/or tempered martensite are generated, resulting in insufficient ductility. For this reason, the coiling temperature is (Bs-250)°C or more, preferably (Bs-230)°C or more, and more preferably (Bs-200)°C or more. On the other hand, if the coiling temperature is higher than (Bs-100)°C, granular bainite is excessively generated, and therefore excellent fatigue strength cannot be obtained. For this reason, the coiling temperature is (Bs-100)°C or less, preferably (Bs-120)°C or less, and more preferably (Bs-150)°C or less.

(6) Cooling (second cooling)
Next, the steel is cooled to (Bs-400) ° C. or less at an average cooling rate of 1 ° C./s or less (second cooling). If the average cooling rate from the coiling temperature to (Bs-400) ° C. or less exceeds 1 ° C./s, lower bainite and/or tempered martensite will be generated, resulting in insufficient ductility. For this reason, the average cooling rate from the coiling temperature to (Bs-400) ° C. or less is 1 ° C./s or less, preferably 0.8 ° C./s or less, and more preferably 0.5 ° C./s or less. There is no particular reason for limiting the lower limit of the average cooling rate, but from the viewpoint of manufacturing efficiency and coarsening of carbides, 0.001 ° C./s or more is preferable. The cooling can be performed to any temperature below (Bs-400) ° C., but it is preferable to cool to 50 ° C. The cooling can be performed in any form, for example, in the state of a coiled coil.

The high-strength steel plate of the present invention can be manufactured by the above procedure. After coiling and subsequent cooling, the plate can be processed in the usual manner. For example, temper rolling can be performed, and pickling can be performed to remove scale formed on the surface.

Molten steel having the composition shown in Table 1 was melted in a converter and cast at the rotation speed shown in Table 2 by a continuous casting method to produce a steel slab as a steel material. The obtained steel material was heated to the heating temperature shown in Table 2, and then the heated steel material was subjected to hot rolling consisting of rough rolling and finish rolling to produce a hot-rolled steel sheet. The finish rolling end temperature in the hot rolling was as shown in Table 2. Next, the obtained hot-rolled steel sheet was cooled under the conditions of the average cooling rate and cooling stop temperature shown in Table 2 (first cooling). The cooled hot-rolled steel sheet was coiled at the coiling temperature shown in Table 2, and the coiled steel sheet was cooled at the average cooling rate shown in Table 2 (second cooling) to obtain a high-strength steel sheet. After cooling, temper rolling was performed and pickling was performed. Pickling was performed at a temperature of 85°C using a hydrochloric acid aqueous solution with a concentration of 10% by mass.

Test specimens were taken from the obtained high-strength steel plates, and the microstructure and mechanical properties were evaluated using the procedures described below.

<Microstructure>
A test piece for microstructure observation was taken from the obtained high-strength steel plate so that the plate thickness cross section parallel to the rolling direction was the observation surface. The surface of the obtained test piece was polished, and the surface was further corroded using an etching solution (3% nital solution) to reveal the microstructure. Next, the plate thickness 1/4 position was photographed at 5000 times magnification using a scanning electron microscope (SEM) to obtain a SEM image of the microstructure. The obtained SEM image was analyzed by image processing to quantify the area ratio of upper bainite (UB), polygonal ferrite (F), lower bainite and/or tempered martensite (LB + TM). In addition, since it is difficult to distinguish between upper bainite (UB), granular bainite (GB), fresh martensite (FM), and retained austenite phase (γ), they were identified using electron beam reflection diffraction method, and the area ratio and average crystal grain size of each were obtained. The area ratio of each microstructure and the average crystal grain size of upper bainite measured are shown in Table 3. In addition, Table 3 also shows the total area ratio (B) of upper bainite and granular bainite, and the total area ratio (FM+γ) of fresh martensite and retained austenite phase.

<Tensile test>
From the obtained hot-rolled steel sheet, JIS No. 5 tensile test pieces (JIS Z 2201) were taken so that the tensile direction was perpendicular to the rolling direction, and a tensile test was carried out in accordance with the provisions of JIS Z 2241, in which the strain rate is 10 ^-3 /s, to determine the tensile strength and uniform elongation. In the present invention, a tensile strength of 980 MPa or more was considered to be acceptable. Furthermore, a uniform elongation of 6.0% or more was evaluated as excellent ductility.

<Plane bending fatigue test>
From the obtained hot-rolled steel sheet, a test piece having the dimensions and shape shown in FIG. 1 was taken so that the longitudinal direction of the test piece was perpendicular to the rolling direction, and a plane bending fatigue test was carried out in accordance with the provisions of JIS Z 2275. The stress load mode was a stress ratio R = -1 and a frequency f = 25 Hz. The load stress amplitude was changed to six stages, the stress cycle up to break was measured, an S-N curve was obtained, and the fatigue strength (fatigue limit) at 2 x ¹⁰⁶ times was obtained. In the present invention, when the value obtained by dividing the fatigue limit by the tensile strength obtained in the tensile test was 0.50 or more, it was evaluated as excellent fatigue properties.

<Punching property evaluation>
Test pieces (size: t (sheet thickness) x 30 mm (width) x 30 mm (length)) were taken from the obtained hot-rolled steel sheet. A punch was made in the center of the taken test piece by punching with a cylindrical punch of 10 mmφ at three or more clearances including 10% and 20% within the clearance range of 10 to 20% to form a punched hole. The clearance is a percentage [%] of the sheet thickness of the test piece. When no cracks, chips, brittle fracture surfaces, or secondary shear surfaces were visually observed on the end surface of the punched hole, it was evaluated as having excellent punchability (○), and when these were confirmed, it was evaluated as ×.

<Hardness measurement>
From the obtained high strength steel plate, a sample for hardness measurement was taken so that the plate thickness cross section parallel to the rolling direction was the hardness measurement cross section, and the hardness (Hv1) at the plate thickness 1/2 position and the hardness (Hv2) at the plate thickness 1/4 position were measured. The Vickers hardness measurement conditions were a load of 100 g and a holding time of 10 s. The measurement points were five points in the ±1/20 plate thickness region at measurement intervals of 250 μm or more, and the average was calculated.

Next, the difference between the obtained Hv1 and Hv2 (|Hv1-Hv2|) was calculated as a percentage relative to 0.3 times the tensile strength obtained in the tensile test and evaluated.

All of the examples of the invention are high-strength steel sheets with tensile strengths of 980 MPa or more and excellent ductility, fatigue strength, and punchability. On the other hand, the comparative examples outside the scope of the present invention do not achieve a tensile strength of 980 MPa or more, or do not achieve excellent ductility, fatigue strength, or punchability.

Claims (3)

In mass percent,
C: 0.04-0.18%,
Si: 0.1 to 2.5%,
Mn: 1.0 to 3.0%,
P: 0.1% or less,
S: 0.01% or less,
Contains Al: 0.010 to 1.0% and N: 0.01% or less;
The remainder is Fe and unavoidable impurities.
The microstructure is mainly composed of bainite phase consisting of upper bainite and granular bainite with an area ratio of 85% or more.
The area ratio of a hard second phase including fresh martensite and/or retained austenite phase is less than 15%;
In the bainite phase, the area ratio of granular bainite to the area ratio of upper bainite is 5% or more and 40% or less,
The difference |Hv1-Hv2| between the hardness Hv1 at 1/2 the sheet thickness position and the hardness Hv2 at 1/4 the sheet thickness position is 30% or less with respect to 0.3TS, where Hv1 and Hv2 are Vickers hardness (Hv), and TS is tensile strength (MPa);
The high-strength steel plate has an average crystal grain size of the bainite phase of 7 μm or less. The high-strength steel plate according to claim 1, further comprising, in addition to the above-mentioned component composition, one or more of groups a to e in mass%.
Group A:
V: 0.005-0.5%,
One or more selected from Ti: 0.005 to 0.2% and Nb: 0.005 to 0.1%;
Group B:
Cu: 0.005 to 0.5%,
Ni: 0.005 to 0.5%,
One or more selected from the group consisting of Cr: 0.005 to 1.0% and Mo: 0.005 to 0.5%;
Group C:
B: 0.0002 to 0.005%,
Group d:
One or two selected from Sb: 0.001 to 0.1% and Sn: 0.001 to 0.1%;
Group e:
Ca: 0.0001-0.005%,
One or more selected from Mg: 0.0001 to 0.005%, and REM: 0.0001 to 0.005%. A method for producing a high strength steel plate according to claim 1 or 2,
A molten steel material having the above-mentioned composition is cast while being rotated in a horizontal plane relative to a mold at a rotation speed of 10 cm/s or more by an induction electromagnetic stirrer;
The obtained steel material is heated to 1150°C or higher,
The heated steel material is roughly rolled into steel plates,
The steel sheet is finish-rolled under the condition of a finish rolling end temperature of (RC-50)°C or more and (RC+100)°C or less,
The steel sheet after the finish rolling is cooled under the conditions of a time from the end of the finish rolling to the start of cooling: 2.0 s or less, an average cooling rate on the surface up to Bs: 30 ° C./s or more, a residence time at a temperature of (Bs-100) ° C. or more and Bs ° C. or less: 3.0 s or more and 10.0 s or less, and a cooling stop temperature: (Bs-250) ° C. or more and (Bs-100) ° C. or less,
The cooled steel sheet is coiled under the condition of a coiling temperature of (Bs-250)°C or more and (Bs-100)°C or less.
A method for producing high strength steel plate, comprising cooling to (Bs-400)°C or lower at an average cooling rate of 1°C/s or less.
Here, RC and Bs are defined by the following equations (1) and (2), respectively.
RC (℃) = 750 + 120 x C + 100 x N + 10 x Mn + 250 x Ti + 5000 x B + 10 x Cr + 50 x Mo + 750 x Nb + 150 x V... (1)
Bs (℃) = 830-270 x C-90 x Mn-70 x Cr-37 x Ni-83 x Mo-20 x Cu... (2)
Here, each element symbol in the above formulas (1) and (2) represents the content (mass%) of each element, and when no element is contained, it is set to 0.

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