JP5076959B2 - Low yield ratio high strength steel sheet with excellent ductile crack initiation characteristics and its manufacturing method - Google Patents

Description

  TECHNICAL FIELD The present invention relates to a structural steel plate suitable for use in buildings, offshore structures, shipbuilding, bridges, line pipes, etc., and in particular, a tensile excellent in ductile crack initiation characteristics required for steel plates used in earthquake-prone areas. The present invention relates to a low yield ratio high strength steel plate having a strength of 550 MPa or more and a method for producing the same.

  In recent years, steel materials used in the fields of buildings, offshore structures, shipbuilding, bridges, line pipes, etc. have improved safety and operational efficiency (for example, increased transport gas pressure in pipelines), Higher strength is being promoted for the purpose of reducing the total cost by reducing the amount of steel used. Moreover, since the area where the above steel materials are used has expanded to the harsh areas of the natural environment, for example, structural steel materials used in earthquake-prone areas are different from the conventional required performance. Therefore, excellent plastic deformability and ductile fracture resistance have been demanded.

  From such a situation, in Patent Documents 1 to 3, steel materials having improved plastic deformability by reducing the yield ratio, which is the ratio of yield stress to tensile strength, are disclosed in Patent Documents 4 and 5. Similarly, high deformability steel materials having excellent buckling resistance properties by reducing the yield ratio have been proposed. However, even if the steel material has a low yield ratio and excellent deformability, if a ductile crack occurs from a stress-concentrated part such as a defect, and this progresses, before the plastic deformability is exerted In addition, cracks may propagate over long distances and cause unstable ductile fracture.

Therefore, development of high-strength and high-deformability steel materials has been carried out for the purpose of preventing unstable ductile fracture. For example, Patent Document 6 discloses a method for producing high-tensile steel in which absorbed energy is increased by making the metal structure a fine-grain ferrite-based structure, and Patent Document 7 describes that the metal structure is composed of ferrite and martensite. A high-strength steel pipe with improved stopping performance of unstable ductile fracture has been proposed.
JP-A-55-119152 JP 63-223123 A Japanese Patent Laid-Open No. 03-115524 Japanese Patent Laid-Open No. 10-330885 Japanese Unexamined Patent Publication No. 2000-178689 JP 2002-105534 A JP 2004-197191 A

  The technologies described in Patent Documents 6 and 7 relate to a steel material with improved ductile crack propagation characteristics in unstable ductile fracture. However, for example, in a gas pipeline operated at high pressure, once a crack has occurred, it may be difficult to stop the propagation of the crack, and as a result, there is a concern that even local destruction may cause serious damage. Has been. For this reason, steel materials used in gas pipelines not only suppress the propagation of ductile cracks, but even if damage occurs, they do not lead to crack generation, or even if cracks occur, gas leakage is minimized. It is desirable that it can be stopped to the limit.

  In order to meet the above requirements, it is necessary to be able to surely prevent the damage caused by the defects inherent in the steel material itself or external factors, or the occurrence of cracks from the reduced thickness portion due to corrosion. However, the present situation is that there is no low yield ratio high strength steel plate with high ductile crack initiation resistance that satisfies the above characteristics.

  Accordingly, an object of the present invention is to provide a low yield ratio high strength steel sheet that is excellent in ductile crack initiation characteristics required for steel materials used in earthquake-prone areas, and to produce the steel sheet at low cost and efficiently. It is to propose an advantageous manufacturing method that can be performed.

  In order to solve the above-mentioned problems, the inventors have conducted intensive studies on the ductile fracture behavior of a high-strength steel sheet having a notch, assuming that damage due to external factors exists. As a result, the cooling is temporarily stopped in the accelerated cooling process, and then immediately reheated to obtain a three-phase structure consisting of a ferrite phase, a bainite phase, and an island martensite phase, thereby optimizing the microstructure structure. Further, it has been found that by controlling the contents of Ca, O and S within an appropriate range, the occurrence of a ductile crack from a notch can be suppressed without reducing the deformability. The invention has been completed.

That is, the present invention is C: 0.03-0.1 mass%, Si: 0.01-1 mass%, Mn: 1.2-2.5 mass%, S: 0.002 mass% or less, Al: 0.01 -0.07mass%, Ca: 0.001-0.003mass%, O: 0.003mass% or less is contained, Ca, S, and O are following (1) Formula;
0.8 ≦ (1-130 × O) × Ca / (1.25 × S) ≦ 2.0 (1)
Here, the element symbol in the formula (1) is the mass% of each element.
In which the balance is composed of Fe and inevitable impurities, the metal structure is composed of a three-phase structure of ferrite phase, bainite phase and island martensite phase, and the phase fraction of island martensite Is a low-yield-ratio high-strength steel sheet having a structure with an average aspect ratio of 8 or less and a yield ratio of 0.80 or less and having excellent ductile crack initiation characteristics.

  In addition to the above component composition, the low yield ratio high-strength steel sheet of the present invention further includes Nb: 0.005-0.1 mass%, V: 0.005-0.1 mass%, and Ti: 0.005-0.1 mass. 1 type or 2 types or more chosen from% are characterized by the above-mentioned.

  In addition to the above component composition, the low yield ratio high-strength steel sheet of the present invention further includes Cu: 0.01 to 0.5 mass%, Ni: 0.05 to 0.5 mass%, Cr: 0.01 to 0. It is characterized by containing 1 type (s) or 2 or more types chosen from 0.5 mass% and Mo: 0.01-0.5 mass%.

In addition, the present invention is a method in which a steel slab having any of the above-described component compositions is heated to 1000 to 1300 ° C., and then hot-rolled with a rolling end temperature equal to or higher than the Ar ₃ transformation point, and a cooling rate of 5 ° C./s. This is excellent in ductile cracking characteristics in which accelerated cooling is performed to a temperature of 450 ° C. to 650 ° C., and then immediately reheated to 550 to 700 ° C. above the accelerated cooling stop temperature at a heating rate of 0.5 ° C./s or higher. A method for producing high yield strength steel sheets is proposed.

  ADVANTAGE OF THE INVENTION According to this invention, even when it receives big plastic deformation by an earthquake etc., the low yield ratio high strength steel plate which does not generate | occur | produce a ductile crack can be provided. Therefore, the steel plate of the present invention is suitable for use in buildings, offshore structures, shipbuilding, bridges, line pipes, and the like that have high safety requirements.

  In the process of accelerated cooling of the steel sheet after hot rolling, the present invention stops cooling in the middle of the bainite transformation, that is, in a temperature region equal to or higher than the bainite transformation finish temperature where untransformed austenite still exists, and then the temperature Immediately after heating, and by concentrating C in untransformed austenite, in the subsequent cooling process, an island-like martensite phase (hereinafter referred to as “MA phase”) which is a hard phase in the mixed phase of ferrite phase and bainite phase. It is characterized in that a three-phase structure composed of a ferrite phase that is a soft phase, a bainite phase, and an MA phase that is a hard phase is obtained. The steel sheet of the present invention has a low yield ratio that has a tensile strength of 550 MPa or more and a low yield ratio of 0.80 or less and is excellent in ductile crack initiation characteristics by having the above three-phase structure. A high-strength steel sheet can be obtained. The MA phase can be easily distinguished from other phases by electrolytic etching after etching with a 3 mass% nital solution (nitral alcohol solution).

Next, an experiment that has become an opportunity to develop the low yield ratio high strength steel sheet of the present invention will be described.
From various steel materials (thick steel plates) that have a three-phase structure of ferrite phase, bainite phase, and MA phase, and have different MA phase fractions, as shown in FIG. A round bar test piece having an annular notch with a radius of 0.25 mm was sampled and subjected to a tensile test. The average strain between the gauge points when a crack generated from the notch bottom of the round bar test piece was first confirmed was “ The ductile crack initiation strain was defined, and the resistance to ductile crack initiation was evaluated based on the magnitude of this strain. At the same time, the microstructure of the steel was investigated and the relationship with the ductile crack initiation characteristics was investigated.

  FIG. 2 is a graph showing the relationship between the phase fraction of the MA phase of a steel sheet having a three-phase structure, the yield ratio, and the ductile crack initiation strain. From FIG. 2, the yield ratio shows a low value of 0.80 or less over a wide range where the MA phase fraction is 3% or more, but the ductile crack initiation strain decreases as the MA phase fraction increases. However, when it exceeds 15%, it turns out that it falls rapidly.

  The reason why the ductile crack initiation strain is high up to a range where the phase fraction of the MA phase is high in the steel plate having a three-phase structure is considered as follows. When a steel sheet having a three-phase structure of a ferrite phase, a bainite phase, and an MA phase is subjected to processing, the ferrite phase generated by reheating treatment and the tempered bainite phase are softer than the MA phase. And the bainite phase begins to deform from the beginning of processing. On the other hand, since the tensile strength is ensured by the hard MA phase, the ratio of yield stress to tensile strength, that is, the yield ratio (= yield stress / tensile strength) is lowered, so that the work hardening index is increased. , Deformability is improved.

  Therefore, the inventors further examined the above results, and after appropriately controlling the component composition and production conditions, further controlling the phase fraction and average aspect ratio of the MA phase to an appropriate range. The present inventors have found that a low-yield-ratio high-strength steel sheet having excellent ductile crack initiation characteristics can be obtained, and the present invention has been completed.

Next, the component composition that the steel sheet of the present invention should have will be described.
C: 0.03-0.1 mass%
C is an element necessary for promoting the formation of the MA phase and increasing the strength of the steel sheet. If the addition is less than 0.03 mass%, the MA phase fraction is low and the desired strength cannot be obtained. On the other hand, if the addition exceeds 0.1 mass%, the weldability decreases. Therefore, C is set to a range of 0.03 to 0.1 mass%.

Si: 0.01-1 mass%
Si is an element added as a deoxidizer and to increase the strength of steel. If it is less than 0.01 mass%, the deoxidation effect is not sufficient, while addition exceeding 1 mass% decreases toughness and weldability. Therefore, Si is set to a range of 0.01 to 1 mass%.

Mn: 1.2 to 2.5 mass%
Mn is added to increase the strength and toughness of the steel and to increase the hardenability and promote the formation of the MA phase. However, if the addition is less than 1.2 mass%, the effect is not sufficient, while the addition exceeding 2.5 mass% decreases the weldability. Therefore, Mn is set to a range of 1.2 to 2.5 mass%.

S: 0.002 mass% or less S is an element that is inevitably mixed as an impurity, and generally exists as a sulfide-based inclusion in steel and serves as a starting point for void generation during deformation. Therefore, in order to prevent the occurrence of ductile cracks, the S content must be strictly regulated. However, if it is 0.002 mass% or less, the above-mentioned adverse effect is small. Therefore, S has an upper limit of 0.002 mass%. Preferably it is 0.001 mass% or less.

Al: 0.01-0.07 mass%
Al is added as a deoxidizing agent in the steel making process, as is the case with Si. However, if the addition is less than 0.01 mass%, the deoxidation effect is not sufficient, while the addition exceeding 0.07 mass% is an oxide type. Increasing amounts of inclusions reduce toughness. Therefore, Al is set to a range of 0.01 to 0.07 mass%.

Ca: 0.001 to 0.003 mass%
Ca is an effective element for controlling the form of sulfide inclusions and improving ductility, but the effect cannot be obtained with addition of less than 0.001 mass%. On the other hand, even if added over 0.003 mass%, the above effect is saturated. Moreover, toughness falls by the fall of a cleanliness or the production | generation of coarse CaO, and it becomes the cause of the nozzle blockade of a ladle, and inhibits productivity. Therefore, Ca is set to a range of 0.001 to 0.003 mass%.

O: 0.003 mass% or less O is an unavoidable impurity, generates coarse oxide inclusions, and adversely affects ductility and toughness. Therefore, in the present invention, O is limited to 0.003 mass% or less.

0.8 ≦ ((1-130 × O) × Ca) / (1.25 × S) ≦ 2.0
In the steel of the present invention, in addition to satisfying the above component composition, it is necessary that Ca, O, and S further satisfy the following formula (1).
0.8 ≦ ((1-130 × O) × Ca) / (1.25 × S) ≦ 2.0 (1)
(However, each element symbol in the above formula is the content of each element (mass%))
This formula (1) suppresses the formation of coarse sulfide inclusions that are the starting point of ductile crack generation, and controls the sulfide form so as not to affect the generation of ductile cracks. After controlling to 001 to 0.003 mass%, S: 0.002 mass% or less, the effective Ca amount (Ca *) excluding the Ca content consumed by the generation of CaO from the Ca content is obtained by experiments. The following regression formula (2):
Ca * = (1-130 × O) × Ca (2)
(However, each element symbol in the above formula is the content of each element (mass%))
Then, based on the stoichiometric ratio of Ca and S, the effective Ca to S atomic ratio (Ca * / (1.25 × S)) is in the range of 0.8 to 2.0. It is shown that it is necessary to control.
That is, when Ca * / (1.25 × S) is less than 0.8, S cannot be sufficiently fixed as CaS, so that coarse MnS is generated and ductile crack resistance is deteriorated. When /(1.25×S) exceeds 2.0, coarse Ca-based oxides and sulfides are generated and the ductile crack generation characteristics are lowered. Therefore, Ca, S, and O are the above (1 ) It is necessary to contain and satisfy the formula.

The steel plate of the present invention has Cu: 0.01 to 0.5 mass%, Ni: 0.01 to 0.5 mass%, Cr: 0.01 to 0.5 mass%, and Mo: 0.0. One or more selected from 01 to 0.5 mass% can be contained.
Cu, Ni, Cr and Mo are elements that increase the strength and toughness of the steel, further improve the hardenability and promote the formation of the MA phase, and can be added according to the required strength. For each element, a sufficient effect cannot be obtained if it is added in an amount of less than 0.01 mass%. On the other hand, if it exceeds 0.5 mass%, weldability is deteriorated. Therefore, when adding these elements, it is preferable to add in the range of 0.01-0.5 mass%, respectively.

Moreover, the steel plate of this invention is chosen from Nb: 0.005-0.1mass%, V: 0.005-0.1mass%, and Ti: 0.005-0.1mass% other than the said component. 1 type, or 2 or more types can be contained.
Nb, V, and Ti are elements that increase the strength and toughness of the steel sheet, and can be added according to the required strength. For each element, a sufficient effect cannot be obtained if it is added in an amount of less than 0.005 mass%. On the other hand, if it exceeds 0.1 mass%, the toughness of the welded portion is lowered. Therefore, when adding these elements, it is preferable to add in the range of 0.005-0.1 mass%, respectively.

  In the steel sheet of the present invention, the balance other than the above components consists of Fe and inevitable impurities. However, even if components other than those described above are added in an amount that is normally recognized as an inevitable impurity, it is allowed as long as the effects of the present invention are not impaired.

Next, the microstructure of the steel sheet of the present invention will be described.
Microstructure: Three-phase structure composed of ferrite phase, bainite phase, and MA phase The steel sheet of the present invention needs to have a microstructure composed of three phases of a soft ferrite phase, a bainite phase, and a hard MA phase. The microstructure consisting of the above three phases is such that in the accelerated cooling process after hot rolling, the cooling is temporarily stopped during the bainite transformation, that is, in the temperature range where untransformed austenite exists, and then immediately reheated, from the untransformed austenite to the ferrite. Can be obtained without reducing the production efficiency by concentrating C in the untransformed austenite and generating the MA phase in the subsequent cooling process.

  FIG. 3 is a structural photograph of the microstructure of the steel sheet of the present invention containing C: 0.05 mass% C, Si: 0.2 mass%, Mn: 1.8 mass% observed with a scanning electron microscope (SEM). It can be confirmed that the MA phase is uniformly dispersed in the mixed phase of the ferrite phase and the bainite phase.

MA phase fraction: 3-15%
The phase fraction of the MA phase needs to be in the range of 3 to 15% in order to obtain a steel sheet having a low yield ratio and excellent ductile crack initiation characteristics. As can be seen from FIG. 2 described above, when the MA phase fraction is less than 3%, the yield ratio is high, whereas when it exceeds 15%, a ductile crack is easily generated. In order to obtain a more excellent low yield ratio and ductile crack initiation characteristics, the MA phase fraction is preferably in the range of 5 to 12%. Here, the phase fraction of the MA phase is the volume fraction (%) of the MA phase with respect to the entire structure, and the volume fraction can be replaced with the area fraction (%) of the cross section of the steel sheet.

  In the steel sheet of the present invention, the remainder other than the MA phase is mainly a ferrite phase and a bainite phase, but from the viewpoint of securing the steel sheet strength, the phase fraction of the bainite phase is preferably 10% or more. In addition to the three phases of ferrite phase, bainite phase and MA phase, heterogeneous phases such as pearlite phase and cementite phase can exist. However, when these different phases are present, the strength decreases, so that the smaller the phase fraction of these different phases, the better. However, if the total of these phase fractions is 5% or less, it does not adversely affect the function and effect of the present invention, and thus is allowed.

Average aspect ratio of MA phase: 8 or less The average aspect ratio of the MA phase needs to be 8 or less. When the average aspect ratio of the MA phase exceeds 8, strain is concentrated at the interface with the ferrite phase and bainite phase near the tip of the elongated MA phase during deformation, and becomes a starting point of a ductile crack. This is because the crack generation characteristics deteriorate. Here, the aspect ratio of the MA phase is defined as the ratio of the major axis to the minor axis (major axis / minor axis) of the elongated MA phase as shown in FIG. The average aspect ratio is an average value of the aspect ratios of 300 MA phases.

Next, the mechanical characteristics of the steel sheet of the present invention will be described.
Yield ratio: 0.80 or less The present invention is intended for a high-strength steel sheet having a tensile strength of 550 MPa or more, but the steel sheet needs to have a yield ratio of 0.80 or less. When the yield ratio exceeds 0.80, not only the deformability of the steel material decreases, but also the strain concentration near the defect portion increases, resulting in localized plastic deformation in the cross section having the defect, This is because ductile cracks easily occur. Preferably, the yield ratio is 0.75 or less.

Ductile crack initiation strain: 2.5% or more In addition to having a low yield ratio, the steel sheet of the present invention needs to be resistant to ductile cracks. The resistance to ductile cracking can be evaluated by ductile crack initiation strain, and the steel sheet of the present invention is required to have the ductile crack initiation strain of 2.5% or more. Here, the above-described ductile crack initiation strain refers to the unloading after applying various amounts of tensile strain to the round bar test piece having an annular notch with a notch bottom radius of 0.25 mm shown in FIG. When observing the structure of the notch bottom cross section of the test piece, it is the average strain between the gauge points when the occurrence of a ductile crack was first observed. It should be noted that the excellent ductile crack initiation characteristic with a ductile crack initiation strain of 2.5% or more can be realized by making the structure of the steel sheet having the above-described component composition the above-described three-phase structure.

Next, the manufacturing method of the low yield ratio high strength steel plate according to the present invention will be described.
Slab heating temperature: 1000-1300 ° C
The steel sheet of the present invention is a steel slab obtained by melting a steel having a composition suitable for the above-described present invention by a normal method through secondary refining such as a converter, an electric furnace or further vacuum degassing. Thereafter, the steel slab is heated in a heating furnace and hot rolled. Here, the heating temperature of the steel slab needs to be in the range of 1000 to 1300 ° C. When the heating temperature is less than 1000 ° C, the solid solution of the carbide becomes insufficient, so that sufficient strength cannot be obtained. On the other hand, heating exceeding 1300 ° C causes the steel sheet structure after rolling to become coarse and toughness is reduced. It is.

Rolling end temperature After the slab is heated, it is hot rolled and processed into a steel sheet. The rolling end temperature in this hot rolling needs to be not less than the Ar ₃ transformation point. If the rolling end temperature is less than the Ar ₃ transformation point, the ferrite transformation rate that occurs during the subsequent cooling decreases, and during reheating after accelerated cooling, the concentration of C into untransformed austenite becomes insufficient, making it difficult to produce the MA phase. Become. Furthermore, when the rolling end temperature is less than the Ar ₃ transformation point, proeutectoid ferrite is precipitated and rolled, and thus the processed ferrite phase remains, so that the ductile crack generation characteristics are deteriorated. In addition, although it does not prescribe | regulate especially about the upper limit of rolling completion temperature, it is preferable that it is 950 degrees C or less in order to roll in a non-recrystallized area and to refine | miniaturize a structure | tissue. Here, the Ar ₃ transformation point varies depending on the component composition of the steel sheet, and can usually be obtained by the following equation (3).
Ar ₃ transformation point (° C.) = 910-310 × C-80 × Mn-20 × Cu-15 × Cr-55 × Ni-80 × Mo (2)
Here, the element symbol of the above formula is the content (mass%) of each element.

Accelerated cooling In the present invention, untransformed austenite enriched with C is transformed into the MA phase by subsequent cooling by reheating after accelerated cooling following hot rolling described later. For this purpose, the steel sheet after hot rolling is cooled from a rolling end temperature to a temperature of 450 to 650 ° C. that is not lower than the bainite transformation end temperature ( _Bf point) where untransformed austenite still exists. It is necessary to perform accelerated cooling at a temperature of ° C / s or higher and stop cooling in that temperature range.

  If the cooling stop temperature is less than 450 ° C., the bainite transformation is completed, and therefore, the MA phase is not sufficiently generated during cooling after reheating. Moreover, even if untransformed austenite remains in part and the MA phase is formed, since it precipitates in a needle shape between the grain boundaries or laths of bainite, the aspect ratio of the MA phase increases and ductile crack initiation characteristics Decreases. On the other hand, if the cooling stop temperature exceeds 650 ° C., C is consumed in the pearlite that precipitates during cooling, and C does not concentrate in the untransformed austenite, making it difficult to produce the MA phase. Therefore, the accelerated cooling stop temperature is in the range of 450 to 650 ° C. In order to further promote the formation of the MA phase, 500 to 650 ° C. is preferable.

  Moreover, the reason why the accelerated cooling rate is set to 5 ° C./s or more is that when it is less than 5 ° C./s, a large amount of pearlite phase is generated during cooling, and high strength cannot be obtained by the bainite phase. When it is desired to fully utilize transformation strengthening by the bainite phase, the accelerated cooling rate after the rolling is preferably 10 ° C./s or more. There is no restriction | limiting in particular about the method of accelerated cooling, Arbitrary facilities can be used by a manufacturing process.

Reheating treatment Next, in the production method of the present invention, it is important to immediately perform reheating after accelerating cooling is stopped in the temperature range above the _Bf point where untransformed austenite exists. It is necessary to reheat from the stop temperature to a temperature of 550 to 700 ° C. that is higher than the bainite transformation end temperature (B _f point) at a temperature increase rate of 0.5 ° C./s or more. If the rate of temperature rise is less than 0.5 ° C./s, pearlite transformation occurs during heating, so that a sufficient amount of MA phase cannot be obtained, and it takes a long time to reach the desired reheating temperature. Efficiency deteriorates.

Further, when the reheating temperature is equal to or lower than the _Bf point, the bainite transformation is completed and untransformed austenite does not exist. Furthermore, when the reheating temperature is less than 550 ° C., C concentration to untransformed austenite becomes insufficient, and the MA phase is hardly generated. On the other hand, if the reheating temperature exceeds 700 ° C., sufficient strength cannot be obtained due to softening of the bainite phase. In order to reliably concentrate C into untransformed austenite, it is preferable to raise the temperature by 50 ° C. or more from the reheating start temperature (accelerated cooling stop temperature).

  In addition, at the said reheating temperature, it is not necessary to provide holding time in particular, and even if it cools immediately after reheating, sufficient MA phase can be obtained. However, in order to further promote the concentration of C and secure a sufficient amount of MA phase fraction, it is preferable to provide a holding time of 30 minutes or less. When heated and held for more than 30 minutes, dislocation recovery in the bainite phase occurs, and the strength may decrease. Although cooling after reheating is not particularly limited, the FA phase can be sufficiently generated even by air cooling (cooling by standing).

Next, equipment for manufacturing the steel sheet of the present invention will be described.
A heating device for performing reheating after accelerated cooling is preferably installed on the downstream side of a cooling facility that performs accelerated cooling. As the heating device, a gas combustion furnace or induction heating device capable of rapid heating of the steel plate can be used, but the induction heating device can heat the cooled steel plate more quickly than the gas combustion furnace, Temperature control is also easy and the cost is relatively low, which is preferable. In addition, arranging a plurality of induction heating devices in series can be done by simply changing the number of induction heating devices to be used and the power supply even if there is a large change in line speed or type / size of steel sheet. It is preferable because the heating rate and the reheating temperature can be easily controlled.

  FIG. 4 shows an example of a rolling line equipment row for producing the steel plate of the present invention. In the rolling line 1, a hot rolling mill 3, an acceleration cooling device 4, an induction heating device 5, and a hot leveler 6 are arranged from upstream to downstream, and the induction heating device 5 is downstream of the acceleration cooling device 4. Are installed on the same line. By adopting this arrangement, the reheating treatment can be performed quickly after completion of the accelerated cooling, so that the heating treatment can be performed without excessively reducing the steel plate temperature after the rolling cooling. In addition, when using a gas combustion furnace as a heating device, it may be installed on the same line as the rolling equipment as described above, but after installing the heating furnace on the machine side and reheating the steel plate, It is good also as a structure which returns on a line again.

  According to the present invention described above, a steel slab having the above component composition is applied to the steel plate by applying the above manufacturing method, and the microstructure is imparted to the steel slab even if it is subjected to large plastic deformation caused by an earthquake or the like. Further, it is possible to produce a steel plate having a low yield ratio and a high strength that is less likely to generate a ductile crack at a low cost and with a high efficiency.

  The steels A to M having the composition shown in Table 1 are melted by a conventional steelmaking process to form steel slabs by a continuous casting method. These steel slabs are heated, hot-rolled, and after rolling, Immediately after cooling to a predetermined cooling stop temperature using a water cooling type accelerated cooling facility, reheating using an induction heating furnace or a gas combustion furnace and air cooling, a No. 15 having a plate thickness of 15 mm was obtained. 1-18 thick steel plates were obtained. The details of the manufacturing conditions for each steel sheet are shown in Table 2. As shown in FIG. 4, the reheating apparatus is installed on the same line as the accelerated cooling equipment as shown in FIG. 4, and the gas combustion furnace is installed close to the line after the accelerated cooling. Using.

For each steel plate produced as described above, the microstructure near the central portion of the plate thickness was observed, a structure photograph of 10 fields of view was obtained, this was subjected to image analysis, the phase fraction of the MA phase and the average aspect of the MA phase The ratio was determined. The average aspect ratio is an average value of 300 MA phases.
Further, from each of the steel plates, a round bar test piece having a parallel portion of 6 mmφ × 30 mm was collected and subjected to a tensile test to measure the tensile strength and the yield ratio.
In addition, the ductile crack initiation strain is a round bar test piece having an annular notch having a parallel portion of 10 mmφ, a score interval of 26 mm, and a notch bottom radius of 0.25 mm as shown in FIG. Unload the sample after applying various amounts of tensile strain, observe the cross-sectional structure of the notch bottom, and determine the average strain between the gauge points when a ductile crack is first observed. Obtained as the generated strain.

  Table 2 shows the measurement results of the microstructure structure, the phase fraction of the MA phase, the average aspect ratio of the MA phase, the tensile strength, the yield ratio, and the ductile crack initiation strain as well as the production conditions. From Table 2, the steel sheets of the examples of the present invention all have a three-phase structure of ferrite phase, bainite phase and MA phase, yield ratio is 0.80 or less, and ductile crack initiation strain is 2.5% or more. It can be seen that it has excellent deformability and ductile crack initiation characteristics. On the other hand, since either the component composition or the manufacturing conditions of the steel plate of the comparative example is out of the scope of the present invention, a three-phase structure of the ferrite phase, the bainite phase, and the MA phase cannot be obtained, or the MA phase As a result, the yield ratio is high or the ductile crack initiation strain is small, and the desired characteristics of the present invention are not obtained.

It is a figure explaining the round bar test piece which has the annular notch used for the measurement of a ductile crack generating distortion. It is a graph which shows the influence which the MA phase fraction has on the yield ratio and ductile crack initiation strain in the three-phase structure steel of ferrite phase, bainite phase and MA phase. It is the photograph which observed the microstructure of the steel plate of the present invention with the scanning electron microscope. It is the schematic explaining an example of the manufacturing line which manufactures the steel plate of this invention.

Explanation of symbols

1: Rolling line 2: Steel plate 3: Hot rolling mill 4: Accelerated cooling device 5: Induction heating device 6: Hot leveler

Claims (4)

C: 0.03-0.1 mass%, Si: 0.01-1 mass%, Mn: 1.2-2.5 mass%, S: 0.002 mass% or less, Al: 0.01-0.07 mass%, Component: Ca: 0.001 to 0.003 mass%, O: 0.003 mass% or less, Ca, S and O satisfy the following formula (1), the balance being Fe and inevitable impurities The metal structure is composed of a three-phase structure of a ferrite phase, a bainite phase, and an island-like martensite phase, and the island-like martensite has a phase fraction of 3 to 15% and an average aspect ratio of 8 or less. A low yield ratio high strength steel sheet with a yield ratio of 0.80 or less.
0.8 ≦ (1-130 × O) × Ca / (1.25 × S) ≦ 2.0 (1)
Here, the element symbol in the formula (1) is the mass% of each element. In addition to the above component composition, Nb: 0.005 to 0.1 mass%, V: 0.005 to 0.1 mass%, and Ti: 0.005 to 0.1 mass%, or one or two selected The low yield ratio high-strength steel sheet according to claim 1, comprising the above. In addition to the above component composition, Cu: 0.01 to 0.5 mass%, Ni: 0.05 to 0.5 mass%, Cr: 0.01 to 0.5 mass%, and Mo: 0.01 to 0.5 mass The low yield ratio high-strength steel sheet according to claim 1, wherein the steel sheet contains one or more selected from%. A steel slab having the component composition according to any one of claims 1 to 3 is heated to 1000 to 1300 ° C, and then hot-rolled at a rolling end temperature of Ar ₃ transformation point or higher, at a cooling rate of 5 ° C / s or higher. A method for producing a low-yield-ratio high-strength steel sheet that is acceleratedly cooled to a temperature of 450 ° C. to 650 ° C., and then immediately reheated to a temperature of 550 to 700 ° C. that is higher than the accelerated cooling stop temperature at a heating rate of 0.5 ° C./s or higher.

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