JP2005171321A - Ultrahigh strength steel sheet having excellent formability and bending workability, and its production method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide an ultrahigh strength steel sheet in which, particularly, tensile strength of ≥980 MPa is attained, further, formability expressed by stretch and stretch flange formability is excellent, and further, bending workability is excellent, and which is suitable for the forming of an automobile component, and to provide its production method. <P>SOLUTION: The ultrahigh strength steel sheet has a composition comprising, by mass, 0.12 to 0.15% C, 1.0 to 1.5% Si, 2.0 to 2.5% Mn, 0.002 to 0.01% N, ≤0.04% P, ≤0.005% S and ≤0.05% Al, and the balance Fe with inevitable impurities, and has a composite structure of a ferritic phase and a low temperature transformation production phase, wherein the average crystal grain size of the ferritic phase is ≤7 μm, and the volume fraction of the low temperature transformation production phase is 40 to 60%. Also, the Vickers hardness HV(F) of the ferritic phase and the Vickers hardness HV(M) of the low temperature transformation production phase satisfy the relation of HV(M)-HV(F)≤350. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention mainly has a tensile strength TS of 980 MPa or more suitable for use in automobile structural parts after being molded, and is excellent in formability represented by elongation and stretch flange characteristics, and excellent bending workability. The present invention relates to a high-strength steel plate and a manufacturing method thereof.

  In recent years, as social demands for improving safety and weight reduction of automobiles have become stricter, automotive structural members have been replaced with conventional mild steel sheets and high-strength cold-rolled steel sheets with a tensile strength TS of 390 to 440 MPa class. Therefore, ultra-high strength steel sheets having a tensile strength TS of 980 MPa or more are being used rapidly. In addition, steel sheets used for automobile structural members are required to have material characteristics such as stretch characteristics, stretch flange characteristics, and bending workability.

  In general, such an ultra-high-strength cold-rolled steel sheet is a so-called structure-strengthened high-strength steel sheet whose metal structure is composed of a relatively hard low-temperature transformation generation phase such as a martensite phase or a bainite phase. . As a high-strength steel sheet that utilizes such a relatively hard low-temperature transformation generation phase as a strengthening mechanism, a relatively hard low-temperature transformation generation phase is dispersed in a soft ferrite ground to simultaneously improve strength and workability. There is a so-called composite structure steel sheet. With the spread and advance of the continuous annealing technology in recent years, high strength cold-rolled steel sheets having excellent elongation and stretch flange characteristics have been widely used.

  Above all, the continuous annealing equipment with water quenching equipment can greatly increase the cooling rate, so it has excellent characteristics in terms of reducing alloy element addition and controlling the low temperature transformation generation phase, as well as product performance. It is possible to produce a steel plate.

  For example, Patent Document 1 and Patent Document 2 disclose steel sheets having a tensile strength of 980 MPa or more as technologies related to high-strength cold-rolled steel sheets, but it is said that such steel sheets have sufficient workability. hard. Patent Document 3 discloses a method for producing a high-strength cold-rolled steel sheet having excellent local ductility. Such a steel sheet exhibits good elongation characteristics but does not have sufficient stretch flange characteristics. Further, Patent Document 4 discloses a steel plate excellent in stretch flange characteristics, but this steel plate has a tensile strength of less than 780 MPa.

In general, strength and workability have a contradictory relationship, and at present, there is no known document that discloses an ultra-high-strength steel sheet having a tensile strength of 980 MPa or more and excellent formability and bending workability.
JP-A-3-277742 JP-A-4-236741 Japanese Examined Patent Publication No. 7-59726 JP-A-4-350

  The object of the present invention is to propose a novel cold-rolled steel sheet and a method for producing the same for solving such problems of the prior art. More specifically, while achieving a tensile strength of 980 MPa or more. Another object of the present invention is to propose an ultra-high-strength steel sheet suitable for forming automobile parts that are excellent in formability represented by elongation and stretch flange characteristics and also have excellent bending workability, and a method for producing the same.

  In order to achieve the above object, the inventors have conducted extensive experiments from the viewpoints of steel components, production conditions, metal structures, and so forth. As a result, the components are optimized and the manufacturing conditions are controlled within an appropriate range in order to maintain both the stretch properties and stretch flange properties, which are generally contradictory properties, at a high property level and to obtain ultra-high strength. By optimizing the ferrite grain size, the fraction of the low-temperature transformation generation phase, and the hardness, the difference in local deformability can be eliminated and the macro uniform deformation can be achieved without reducing the strength level. As a result, it has been found that the above object is achieved.

  Furthermore, in order to obtain such a microstructure, in the water quenching process, strain is uniformly generated inside the steel plate at a large cooling rate, and the temperature rising rate in the subsequent reheating zone is increased to reduce the internal strain. It has been found that it is necessary to open uniformly, and characteristics superior to conventional heat treatment in which the temperature is raised in the reheating zone and maintained for a long time without rapid heating can be easily and stably obtained.

The present invention has been completed based on such findings, and the gist thereof is as follows.
(1)
C: 0.12-0.15 mass%
Si: 1.0-1.5mass%,
Mn: 2.0-2.5mass%,
N: 0.002 to 0.01 mass%
P: 0.04 mass% or less,
S: 0.005 mass% or less and
Al: 0.05 mass% or less, the balance is composed of Fe and inevitable impurities, is composed of a composite structure of a ferrite phase and a low-temperature transformation generation phase, the average crystal grain size of the ferrite phase is 7 μm or less, The volume fraction of the low temperature transformation phase is 40-60%, and the Vickers hardness HV (F) of the ferrite phase and the Vickers hardness HV (M) of the low temperature transformation phase are HV (M) -HV ( F) An ultra-high-strength steel sheet excellent in formability and bending workability characterized by satisfying the relationship of ≦ 350.

(2) C: 0.12 to 0.15 mass%,
Si: 1.0-1.5mass%,
Mn: 2.0-2.5mass%,
N: 0.002 to 0.01 mass%
P: 0.04 mass% or less,
S: 0.005 mass% or less and
Al: containing less 0.05 mass%, the balance being a steel material having a composition of Fe and inevitable impurities, pickling after hot rolling, after cold rolling, if necessary, Ac ₁ transformation point or more 850 After soaking in a temperature range of ℃ or less, quenching treatment is performed by quenching at a cooling rate of 100 ℃ / s or more from the rapid cooling start temperature to at least 200 ℃, and then increasing by at least 100 ℃ from the quenching stop temperature. Super high with excellent formability and bending workability, characterized by heating at a heating rate of ℃ / s or higher, soaking at 300-500 ° C, cooling, and then temper rolling A method for producing a strength steel plate.

(3) The method for producing an ultra-high strength steel sheet having excellent formability and bending workability as described in (2) above, wherein the heating in the tempering treatment is performed by an induction heating method.

  According to the present invention, it is possible to produce an ultra-high strength steel sheet that is excellent in formability represented by elongation characteristics and stretch flange characteristics and also has good bending characteristics.

The reason which limited the steel component of the ultra high strength steel plate in this invention is demonstrated below.
・ C: 0.12-0.15 mass%
C is an indispensable element for strengthening steel using the low temperature transformation phase. The C content needs to be 0.12 mass% or more in order to obtain both a tensile strength of 980 MPa or more and a low temperature transformation generation phase of 40% or more. On the other hand, if the C content exceeds 0.15 mass%, the weldability deteriorates remarkably, the C concentration in the austenite increases, the hardness of the low-temperature transformation generation phase becomes too high, and the stretch flange characteristic decreases. Thus, the volume fraction of the low-temperature transformation generation phase exceeds 60%, and the elongation decreases. Therefore, the C content is in the range of 0.12 to 0.15 mass%.

・ Si: 1.0-1.5mass%
Si is an element that contributes to strength improvement and is an indispensable element for obtaining a suitable low-temperature transformation generation phase by promoting two-phase separation of the ferrite phase and austenite phase during annealing. The above effect is not exhibited when the Si content is less than 1.0 mass%. On the other hand, when Si exceeds 1.5 mass%, Si concentrates on the surface during continuous annealing, reacts with a small amount of water vapor present in the atmosphere, and forms Si-based oxide on the surface. The chemical conversion treatment performed as a pre-treatment is significantly deteriorated, and the adhesion to the coating is remarkably lowered. Therefore, the Si content is in the range of 1.0 to 1.5 mass%.

・ Mn: 2.0-2.5mass%
Mn is an element that plays an important role in suppressing ferrite transformation and obtaining a low-temperature transformation generation phase. Mn also lowers the Ar ₃ transformation point, contributes to finer crystal grains, and has the effect of increasing the balance between the elongation characteristics and the elongation flange characteristics at the same strength. The Mn content needs to be 2.0 mass% or more from the viewpoint of securing the tensile strength TS by stably obtaining the low-temperature transformation generation phase. On the other hand, if the Mn content exceeds 2.5 mass%, the formation of a soft ferrite phase is excessively suppressed, and the characteristic balance of strength characteristics (especially tensile strength TS) and ductility (especially elongation characteristics and stretch flange characteristics) decreases. To do. Moreover, it is necessary to make Mn content into the range of 2.0-2.5 mass% also from a viewpoint of controlling the volume fraction of a low temperature transformation production | generation phase in the range of 40-60%. Therefore, the Mn content is in the range of 2.0 to 2.5 mass%.

・ N: 0.002 ~ 0.01mass%
N, like C, is an element that contributes to the strengthening of steel and the formation of a low-temperature transformation generation phase. It is also partly dissolved in ferrite to control the difference in hardness between the ferrite phase hardness HV (F) and the low temperature transformation phase HV (M) HV (M)-HV (F) It is an effective element. The said effect is expressed by containing N of 0.002 mass% or more. Moreover, since making N content less than 0.002 mass% causes the fall of productivity and a cost increase, the minimum of N content was made into 0.002 mass%. Furthermore, since N is present in the steel as AlN, a large amount of N decreases the workability and also causes a decrease in spot weldability and a deterioration in surface properties, so the upper limit of the N content is 0.01 mass%. It was.

  In the present invention, the balance other than the above components is substantially the composition of iron and inevitable impurities. In addition, it is permissible to contain a trace amount of elements other than those described above as long as the effects of the present invention are not affected.

・ P: 0.04 mass% or less P is an element that has a high solid solution strengthening ability and contributes to an increase in strength. However, if the P content exceeds 0.04 mass%, solidification segregation during casting becomes significant, and internal cracks and Workability deterioration tends to occur. For this reason, it is preferable to make P content 0.04 mass% or less.

-S: 0.005 mass% or less Since S becomes a stress concentration source of the stretch flange, the S content is preferably 0.005 mass% or less.

・ Al: 0.05 mass% or less
Al is an element effective for deacidification and carbide formation. However, the inclusion of Al in excess of 0.05 mass% may not only saturate the effect but also cause deterioration of workability and surface properties. For this reason, it is preferable that Al content shall be 0.05 mass%.

Next, the reason for limiting the steel structure will be described.
-Average grain size of ferrite phase: 7 μm or less By making the average crystal grain size of ferrite phase 7 μm or less and making the ferrite crystal grains finer, the elongation and stretch flange characteristics can be improved without reducing the tensile strength TS. It is possible, and because the refinement of ferrite crystal grains makes the steel sheet structure more uniform, it is high as a result of suppressing the progress of cracks at the interface between the ferrite phase and the low-temperature phase in stretch flange molding Stretch flange characteristics can be achieved. When the average crystal grain size of the ferrite phase exceeds 7 μm, the above effect is not observed. Therefore, the average crystal grain size of the ferrite phase is set to 7 μm or less.

-Volume fraction of low temperature transformation phase: 40-60%
When the volume fraction of the low-temperature transformation phase is increased under constant strength (for example, TS = 980 MPa), the elongation tends to decrease and the elongation characteristics tend to decrease, whereas the stretch flange characteristics improve. There is a tendency. This tendency is considered to be caused by the following. That is, as the amount of low-temperature transformation product phase increases, the C concentration in the low-temperature transformation product phase decreases, and the hardness difference HV (M) -HV (F) between the ferrite phase and the low-temperature transformation product phase decreases. In addition, since the occurrence of cracks at the interface between the ferrite phase and the low temperature phase in the stretch flange molding is suppressed, stretch flange characteristics are improved. However, since the work hardening ability of the ferrite phase is lowered, the elongation characteristic is lowered.

  On the other hand, when the low-temperature transformation generation phase decreases, the C concentration in the low-temperature transformation generation phase increases, and the hardness difference HV (M) -HV (F) between the ferrite phase and the low-temperature transformation generation phase increases. Since the crack generation at the phase interface in the molding is promoted, the stretch flange characteristic is deteriorated. However, since the work hardening ability of the ferrite phase is increased, the elongation characteristics are improved.

  The present inventors have found that there is a place where both elongation and stretch flange characteristics are extremely good, which is not found in the prior art, by controlling the volume fraction of the low-temperature transformation generation phase in the range of 40 to 60%. It was. As described above, when the volume fraction of the low temperature transformation generation phase is less than 40%, the elongation characteristics are further improved, but the stretch flange characteristics are remarkably deteriorated, and when it exceeds 60%, the stretch flange characteristics are improved. Elongation characteristics are significantly reduced. Therefore, in the present invention, the volume fraction of the low-temperature transformation generation phase is set in the range of 40 to 60%.

  The “low-temperature transformation generation phase” here means specifically a tempered martensite phase.

-The Vickers hardness HV (F) of the ferrite phase and the Vickers hardness HV (M) of the low-temperature transformation phase satisfy the relationship of HV (M)-HV (F) ≤ 350. Vickers hardness HV of the ferrite phase When the difference HV (M) -HV (F) between Vickers hardness HV (M) of (F) and the low-temperature transformation formation phase exceeds 350, it is not a uniform structure but a low-temperature transformation in a soft phase such as a ferrite phase. A hard phase such as a generated phase is locally present, and deformability is different at the time of strain introduction, resulting in uneven deformation. In other words, the ferrite phase is deformed by the introduction of strain, but the low-temperature transformation generation phase such as the tempered martensite phase is difficult to deform, so that voids are easily generated at the interface of the low-temperature transformation generation phase. Since the progress of cracks easily occurs at the interface during molding, the stretch flange characteristic is deteriorated. Therefore, the difference HV (M) −HV (F) between the Vickers hardness HV (F) of the ferrite phase and the Vickers hardness HV (M) of the low temperature transformation generation phase was set to 350 or less.

  Next, the reason for limiting the production conditions in the method for producing an ultra high strength steel sheet according to the present invention will be described below. The reason for limiting the components of the steel material used in the manufacturing method is substantially the same as the reason for limiting the steel components of the ultra-high-strength steel plate described above, and thus the description thereof is omitted.

First, after heating the steel slab which is the steel raw material adjusted to the said component, hot rolling is performed. The slab heating temperature may be in the range of 1150 to 1300 ° C. The hot rolling conditions are not specified, but the hot rolled plate grain size is low, such as the hot rolling finishing temperature is lower than the Ar ₃ transformation point, or the cooling rate after the hot rolling is 5 ° C / s or less. If it does not grow significantly, there will be no problem. Conversely, after hot rolling is completed, the rapid rolling of 100 to 300 ° C / s within 1 second is used, and this is combined with the application of large reductions in hot finish rolling. With respect to the act of reducing, the effect of the present invention is not hindered. The coiling temperature does not affect the effect of the present invention as long as the temperature is 400 to 650 ° C.

The hot-rolled steel strip manufactured under the above conditions is not less than the Ac ₁ transformation point in order to obtain the desired austenite phase volume fraction in the subsequent continuous annealing line after pickling or further cold rolling. In order to adjust the volume fraction of the austenite phase by soaking in a temperature range of 850 ° C or lower, after slowly cooling to the desired quenching start temperature, at a cooling rate of 100 ° C / s or higher to a temperature of at least 200 ° C Apply quenching treatment for rapid cooling. Here, the reason for limiting the soaking temperature to 850 ° C. or less is to suppress the grain size of ferrite and initial austenite, and when it exceeds 850 ° C., the grain size during soaking becomes coarse. This is because the average crystal grain size of the ferrite phase in the final structure cannot be adjusted to 7 μm or less. Moreover, the reason why the soaking temperature is limited to the Ac ₁ transformation point or more is to generate an austenite phase.

  The cooling rate from the rapid cooling start temperature requires a certain rate or more in order to obtain a low-temperature transformation generation phase such as a martensite phase or a bainite phase. Is also an important point. Here, it is necessary to introduce transformation strain and thermal strain as uniform as possible into the steel plate, and for this purpose, a cooling rate of 100 ° C./s or more is required. Of course, since high strength is required, it is necessary to rapidly cool to at least 200 ° C. This is because if the quenching stop temperature is higher than 200 ° C., a hard martensite phase cannot be obtained. The cooling rate of 300 ° C./s or more is particularly preferable because the introduction of strain becomes remarkable and the structure is finely and uniformly dispersed, and the effect is increased. In addition, cooling in the jet water is preferable in that uniform and rapid cooling is possible and the effects of the present invention can be obtained more efficiently.

  After the rapid cooling, a tempering process is performed in which heating is performed at a temperature increase rate of 20 ° C./s or more until the temperature is increased by at least 100 ° C. from the rapid cooling stop temperature, soaking at 300 to 500 ° C. and then cooling. This tempering treatment coarsens and softens the carbide in the low-temperature transformation product phase generated by rapid cooling at 100 ° C / s or higher, and the difference in hardness between the ferrite phase and the low-temperature transformation product phase HV (M)- The HV (F) is set to 350 or less, and the strain accumulated uniformly in the steel sheet is uniformly released to improve the shape. If the temperature is raised at least 100 ° C. from the quenching stop temperature, the softening of the low temperature transformation phase is insufficient if the temperature is lower than this, and the rate of temperature increase after that is 20 ° C./s or more. If it is less than s, coarsening and softening of the carbide will be insufficient, and strain release will be uneven and the shape will be poor. The tempering temperature is set to 300 ° C. or more from the viewpoint of coarsening of the carbide and softening of the low-temperature transformation generation phase accompanying it. However, when the tempering temperature exceeds 500 ° C., the strength is remarkably reduced and the desired strength cannot be obtained. Therefore, the tempering temperature is 300 to 500 ° C. The cooling after soaking at 300 to 500 ° C. is not particularly limited.

  In addition, when the heating in the tempering process after the rapid cooling is performed by the induction heating method, the coarsening of the carbide is further promoted by the remarkable increase of C diffusion compared with the atmospheric heating, so that the low temperature transformation generation phase is softened. From the viewpoint of

  The temper rolling is preferably performed at an elongation of 0.1 to 1%.

  Other manufacturing conditions are not mentioned in particular, but slab manufacturing method by ingot or continuous casting, continuous hot rolling with hot hot rolled bar connection in hot rolling, and induction in hot rolling process A temperature rise within 200 ° C. using a heater does not affect the effects of the present invention.

  The above description is merely an example of the embodiment of the present invention, and various modifications can be made within the scope of the claims.

Examples will be described below.
First, the inventive steels A to C and the comparative steels D to I having the component compositions shown in Table 1 were steeled out in a converter and made into steel slabs by continuous casting. After heating these steel slabs at 1250 ° C, they were hot-rolled at a finishing temperature of 870 ° C and wound up at 600 ° C to form hot-rolled steel strips. Subsequently, pickling or further cold rolling was performed to prepare an annealed original sheet. In addition, the width | variety of the annealing raw sheet was 1000 mm in any Example, the plate | board thickness of the hot rolled sheet (pickling board) was 1.8-2.0 mm, and the plate | board thickness of the cold rolled sheet was 1.0-1.4 mm. Subsequently, heat treatment was performed in the continuous annealing line under the conditions shown in Table 2. After the heat treatment, pickling was performed, and then temper rolling was performed to obtain a test steel plate. Moreover, when the component in each produced test steel plate was analyzed, it was not substantially different from the component in the steel slab shown in Table 1. The average grain size (μm) of ferrite phase, volume fraction of low-temperature transformation phase (%), and hardness difference between ferrite phase and low-temperature transformation phase HV (M)-HV ( Table 2 shows F). Table 2 also shows the results of evaluating the characteristics of each test steel plate. In addition, the measurement method of the average crystal grain diameter of a ferrite phase, the volume fraction of a low temperature transformation production | generation phase, the hardness of a ferrite phase and a low temperature transformation production phase, and the test method for performing characteristic evaluation are shown below.

(Measuring method)
(I) Average crystal grain size of ferrite phase The average crystal grain size of the ferrite phase is determined by measuring the position of the center of the plate and the center of the plate thickness in order to avoid the formation of a peculiar structure due to segregation of additive elements in the vicinity of the plate thickness center. The area of the ferrite phase and the number of ferrite phases are determined by image analysis based on the 3000 times SEM image, with the vicinity of the 1/4 position corresponding to the midpoint position of the line segment connecting the thick surface position in the thickness direction. It is a value of n = 3 simple average derived and calculated by the quadrature method.

(Ii) Volume fraction of the low-temperature transformation generation phase The volume fraction of the low-temperature transformation generation phase is the second floor by image analysis based on the SEM image of 5000 times, with the measurement position being in the vicinity of the 1/4 thickness position. Adjustment and area ratio were calculated and calculated by simple averaging with n = 5.

(Iii) Difference in hardness between ferrite phase and low-temperature transformation generation phase Difference in hardness between ferrite phase and low-temperature transformation generation phase HV (M) -HV (F) is the hardness of ferrite phase and low-temperature transformation generation phase. The measurement position was obtained by using a Vickers hardness meter near the plate thickness 1/4 surface position, and the value of hardness HV0.003 at a load of 3 g was calculated by simple averaging at n = 5.

(Test method)
(I) Tensile properties Tensile properties (yield stress YS, tensile strength TS, elongation El) were evaluated based on the tensile test method specified in JIS Z 2241 using No. 5 test piece specified in JIS Z 2201. did.

(II) Stretch flange characteristics The stretch flange characteristics were evaluated by the hole expansion ratio λ. The hole expansion ratio lambda, based on the Japan Iron and Steel Federation Standard JFST1001, punched holes initial diameter d _o = 10 mm, when the expanding holes by increasing the cone punch 60 °, cracks through the thickness direction of the plate Then, the punch was stopped and the punched hole diameter d after penetration through the crack was measured and calculated as hole expansion rate (%) = ((d−d _o ) / d _o ) × 100. The stretch flange characteristic indicates that the larger the hole expansion ratio λ, the better.

(III) Bending workability The bending workability was measured by measuring the limit bending radius in the L direction based on the bending test method specified in JIS Z 2248 using No. 3 test piece specified in JIS Z 2204. It was evaluated from. In addition, bending workability shows that it is so excellent that the numerical value of the L direction limit bending radius is small.

  From the evaluation results shown in Table 2, the examples of the present invention have a tensile strength TS of 980 MPa or more, excellent elongation and stretch flange characteristics, and excellent bending workability. On the other hand, the comparative example has a tensile strength TS of 980 MPa or more, but is inferior in any of elongation, stretch flange characteristics and bending workability.

  According to the present invention, an ultra-high-strength steel sheet that achieves a tensile strength of 980 MPa or more, is excellent in formability represented by elongation and stretch flange characteristics, and is suitable for forming automobile parts that are also excellent in bending workability. Manufacturing becomes possible.

Claims (3)

C: 0.12-0.15 mass%
Si: 1.0-1.5mass%,
Mn: 2.0-2.5mass%,
N: 0.002 to 0.01 mass%
P: 0.04 mass% or less,
S: 0.005 mass% or less and
Al: 0.05 mass% or less, the balance is composed of Fe and inevitable impurities, is composed of a composite structure of a ferrite phase and a low-temperature transformation generation phase, the average crystal grain size of the ferrite phase is 7 μm or less, The volume fraction of the low temperature transformation phase is 40-60%, and the Vickers hardness HV (F) of the ferrite phase and the Vickers hardness HV (M) of the low temperature transformation phase are HV (M) -HV ( F) An ultra-high-strength steel sheet excellent in formability and bending workability characterized by satisfying the relationship of ≦ 350. C: 0.12-0.15 mass%
Si: 1.0-1.5mass%,
Mn: 2.0-2.5mass%,
N: 0.002 to 0.01 mass%
P: 0.04 mass% or less,
S: 0.005 mass% or less and
Al: containing less 0.05 mass%, the balance being a steel material having a composition of Fe and inevitable impurities, pickling after hot rolling, after cold rolling, if necessary, Ac ₁ transformation point or more 850 After soaking in a temperature range of ℃ or less, quenching treatment is performed by quenching at a cooling rate of 100 ℃ / s or more from the rapid cooling start temperature to at least 200 ℃, and then increasing by at least 100 ℃ from the quenching stop temperature. Super high with excellent formability and bending workability, characterized by heating at a heating rate of ℃ / s or higher, soaking at 300-500 ° C, cooling, and then temper rolling A method for producing a strength steel plate.   The method for producing an ultra-high strength steel sheet having excellent formability and bending workability according to claim 2, wherein heating in the tempering treatment is performed by an induction heating method.