JP2003253331A - Method for manufacturing high-tensile-strength steel with high toughness and high ductility - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To efficiently manufacture a high-tensile-strength steel with high toughness and high ductility, which has an adequate strength as a steel for a weld structure, is superior in ductility characteristics and low-temperature toughness, and has a high level of safety. <P>SOLUTION: This manufacturing method comprises heating a steel slab having an appropriate chemical composition and a carbon content, to an AC3 transformation temperature or higher but 1,300°C or lower, hot-rolling it so as to include rolling in an unrecrystallized region of austenite with accumulated rolling reduction of at least 30% in a range between 950°C and the Ar3 transformation temperature, then acceleratingly cooling it from the Ar3 transformation temperature or higher to a temperature in which the austenite fraction becomes 20-70%, at 3-100°C/s, stopping accelerated cooling, then keeping the steel temperature within the accelerated-cooling stopping temperature ±100°C, for 10 s-100 s after the above stop of accelerated cooling,,6y performing a combination of one or more of heating, holding, and cooling at a cooling rate of 0.5°C/s or lower, then cooling it, and tempering it as needed, at a heating temperature of 500°C or higher but the AC1 transformation temperature or lower. <P>COPYRIGHT: (C)2003,JPO

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a method for producing high-strength welding tensile steel having excellent toughness. Specifically, the present invention has a tensile strength of 490 MPa to 780 M.
Steel for structures that is used in general structures that are Pa class and require excellent low-temperature toughness with a toughness assurance temperature of 0 to -40 ° C or less and high uniform elongation at the same time from the viewpoint of seismic resistance and the like. It can be used in general for welded steel structures such as marine structures, pressure vessels, shipbuilding, bridges, buildings, line pipes, etc., but it can have both high uniform elongation and toughness. Therefore, it is particularly useful as a steel material for structures such as buildings and bridges that require earthquake resistance. Further, the form of the steel material is not particularly limited, but it is particularly useful for a steel plate used as a structural member and requiring low temperature toughness, particularly a thick plate, a steel pipe material, or a shaped steel.

[0002]

2. Description of the Related Art Recently, from the viewpoint of safety, centering on high-rise buildings, designs have been designed to prevent collapse of the building even when it is deformed assuming an earthquake. In addition, low yield ratio (yield ratio = yield stress / tensile strength) or / and high uniform elongation property are required as properties required for steel materials. By lowering the yield ratio and increasing the uniform elongation, it becomes possible to increase the energy absorption of the steel material when a large deformation is applied. In particular, considering the fracture resistance, it is considered that high uniform elongation is more effective in that local ductile fracture can be suppressed. It is difficult to secure uniform elongation as the strength increases, but in order to improve uniform elongation at the same strength, an appropriate amount of hard phase such as martensite phase is dispersed in ferrite (α) matrix of soft phase. It is known to be effective. Various methods for producing a duplex stainless steel composed of soft ferrite and a hard phase have been proposed in the past, but ferrite (α) + was added during quenching and tempering heat treatment.
As typified by a method of performing an intermediate heat treatment for heating to an austenite (γ) two-phase region (hereinafter, QLT treatment), basically ferrite as a soft phase and bainite or martensite as a hard phase, or both phases The purpose is to mix mixed tissues.

And the overall strength level and yield ratio,
Ductile properties have been controlled by changing the mixing ratio of these phases. Various manufacturing methods for obtaining the mixed structure of the soft phase and the hard phase have hitherto been proposed. For example, in JP-A-53-23817, after reheating and quenching a steel sheet, the Ac1 transformation point and the Ac3 transformation point are set. A method of reheating between points to form two phases of γ and α and then air-cooling is disclosed. Further, JP-A-4-314824 discloses a method of reheating to a two-phase region and then quenching. Is disclosed.
As an on-line manufacturing method without performing reheating treatment, for example, in Japanese Patent Laid-Open No. 63-286517, after hot rolling from a γ region to a two-phase region, Ar3 is used.
A method is disclosed in which the α phase is generated by air cooling to a temperature 20 to 100 ° C. lower than the transformation point, and then rapidly cooled. After the reheating and quenching, the QLT process including the two-phase heat treatment of reheating between the Ac1 transformation point and the Ac3 transformation point to form two phases of γ and α and then air cooling or water cooling is relatively easy to control the structure. However, since the toughness is extremely deteriorated when the two-phase region heat treatment is left as it is, it is indispensable to further perform the tempering treatment below the AC1 transformation point. Therefore, the QLT process has a problem in that the process is complicated and the productivity is greatly reduced. Further, when the tempering treatment is performed below the AC1 transformation point, the high uniform elongation obtained in the two-phase region heat treatment rather deteriorates due to the strength reduction of the hard phase and the precipitation strengthening in the ferrite phase.

[0004]

In steels that undergo a two-phase zone heat treatment to enhance ductility characteristics represented by uniform elongation, it is not possible to achieve both high uniform elongation and toughness depending on the conventional techniques. It was difficult. In addition, there is a problem that the manufacturing cost becomes high due to various processes. Therefore, the present invention is a thermo-mechanical treatment method which does not rely on re-heating treatment, and which can suppress the manufacturing cost of a high-strength steel having a uniform elongation equal to or higher than that of a conventional two-phase heat treatment material and having good toughness. It is an object of the present invention to provide means for manufacturing by.

[0005]

[Means for Solving the Problems] As a result of detailed examination of the structural requirements for achieving a high uniform elongation as well as a low yield ratio by a two-phase organization of a soft phase ferrite and a hard phase, From the viewpoint of the optical microscope, it is preferable that the main constituent phase of the hard phase is martensite, bainite, or a mixed structure of both, and it is further preferable that the hard phase is finely dispersed. As a result of the metallographic analysis of the uniform elongation improving mechanism, fine dispersion of the hard phase causes minute retained austenite to exist in the hard phase, but this causes martensite during plastic deformation. It has been found that a high uniform elongation can be achieved by the work-induced transformation that transforms into. That is, even when compared with the same fraction, fine austenite stabilizes austenite as compared with coarse austenite, and it becomes possible to leave a large amount of retained austenite in the hard phase even after transformation into the hard phase. Note that retained austenite refers to a stable austenite phase that remains untransformed even after the cooling transformation is completed up to room temperature, and has not yet transformed at some stage during the cooling transformation, but up to room temperature. Is different from austenite which transforms to ferrite or hard phase. The austenite existing during transformation is simply referred to as austenite or untransformed austenite.

As the most general means for finely dispersing the hard phase, in the QLT treatment, the structure before L treatment (heat treatment in the two-phase region) is refined to increase the austenite nucleation sites during the heating of L treatment. Can be considered. However, fine dispersion of the hard phase is not easy with a method that does not rely on reheating treatment. For example, in JP-A-63-286517, after hot rolling from the γ region to the two-phase region,
In the method of air-cooling to a temperature 20 to 100 ° C. lower than the Ar3 transformation point to generate an α phase, and then rapid cooling, a hard phase is generated from the remaining austenite after ferrite is generated from the austenite single phase during slow cooling. Therefore, it is essentially difficult to control the size and morphology of the hard phase to make it finer. Further, with this means, it is difficult to reduce the size of both the ferrite and the hard phase, so there is the problem that it is not easy to secure toughness.

According to the present invention, deterioration of toughness is prevented by performing microstructure control only during hot rolling and subsequent cooling process without forming reheat treatment to form fine hard phase and retained austenite. As a result of investigating a new industrial means capable of increasing uniform elongation without accompanying, by controlling the transformation from supercooled austenite during accelerated cooling,
Appropriate two-phase microstructure of ferrite and hard phase as soft phase and refinement of hard phase, which are prerequisites for refinement of transformation structure, low yield ratio and high uniform elongation, and further, as a result, hard We have found a new way to achieve the formation of retained austenite in the phase. The summary is as follows.

(1)% by mass, C: 0.01 to 0.2
%, Si: 0.01 to 1%, Mn: 0.1 to 2%, A
1: 0.001-0.1%, N: 0.001-0.0
1%, P: 0.02% or less as an impurity,
S: 0.01% or less is contained, the carbon equivalent (Ceq.) Represented by the following formula (A) is 0.3 to 0.6%, and the balance Fe
And a steel slab consisting of unavoidable impurities, the AC3 transformation point or higher
After heating to a temperature of 1300 ° C. or lower, at least a start temperature of 950 ° C. or lower, an end temperature of Ar3 transformation point or higher, and a cumulative rolling reduction of 30% or more, after hot rolling including unrecrystallized zone rolling of austenite, Accelerated cooling at a cooling rate of 3 to 100 ° C./s is performed from a temperature of Ar3 transformation point or higher to a temperature at which the austenite fraction becomes 20 to 70%, and after the accelerated cooling is stopped,
The temperature of the steel is accelerated for 10 seconds to 100 seconds after the accelerated cooling is stopped by performing one or two or more combinations of temperature increase, holding, and cooling at a cooling rate of 0.5 ° C./s or less, and the accelerated cooling stop temperature is set. A method for producing a high-toughness / high-ductility high-strength steel, which is maintained within ± 100 ° C and then cooled. Ceq. = C% + Mn% / 6 + (Cu% + Ni%) / 15+ (V% + Mo% + Cr%) / 5 ... (A)

(2) Acceleration cooling stop temperature ±
After maintaining the temperature within 100 ° C, the cooling rate is 3 to 1
The method for producing a high-toughness / high-ductility high-strength steel according to (1), wherein the second accelerated cooling is performed at 00 ° C./s and the stop temperature is 500 ° C. or less. (3) The method for producing a high-toughness / high-ductility high-strength steel according to (1) or (2), further including tempering at a heating temperature of 500 ° C. or higher and an AC1 transformation point or lower.

(4) Further, in mass%, Ni: 0.01
~ 5%, Cu: 0.01-1.5%, Cr: 0.01-
2%, Mo: 0.01 to 2%, W: 0.01 to 2%,
Ti: 0.003 to 0.1%, V: 0.005 to 0.
5%, Nb: 0.003 to 0.1%, Zr: 0.003
.About.0.1%, Ta: 0.005 to 0.2%, B: 0.
The method for producing a high-toughness / high-ductility high-strength steel according to any of (1) to (3) above, which contains one or more of 0002 to 0.005%. (5) Further, in mass%, Mg: 0.0005 to 0.0
1%, Ca: 0.0005 to 0.01%, Y: 0.00
5 to 0.1%, La: 0.001 to 0.1%, Ce:
High toughness according to any one of (1) to (4), which contains one or more of 0.001 to 0.1%.
Manufacturing method of high ductility and high strength steel.

[0011]

BEST MODE FOR CARRYING OUT THE INVENTION Embodiments of the present invention will be described in detail below. The present invention, by optimizing the chemical composition, in the thermomechanical treatment step, by controlling the transformation from supercooled austenite during accelerated cooling, refinement of the transformation structure, low yield ratio and high uniform elongation. Appropriate two-phase organization of ferrite and hard phase as the soft phase which is the premise of the formation, further refinement of the hard phase and, as a result, to achieve the formation of retained austenite in the hard phase, specifically Is
As shown in FIG. 1 showing the outline of the manufacturing method of heating, rolling and cooling of the present invention, a steel slab having an appropriate chemical composition is treated with AC3.
Hot rolling including austenite unrecrystallized zone rolling at a temperature of not lower than the transformation point and not higher than 1300 ° C., a starting temperature of not lower than 950 ° C., an end temperature of not lower than Ar3 and a cumulative reduction of not lower than 30%. After doing, the cooling rate is 3-1
Accelerated cooling of 00 ° C./s is performed from a temperature above the Ar3 transformation point to a temperature at which the austenite fraction is 20 to 70%,
After the accelerated cooling is stopped, the temperature is raised, maintained (heat retention), and the cooling rate is 0.
After cooling at 5 ° C / s or less, one kind or a combination of two or more kinds is performed, and after maintaining the temperature of the steel within the accelerated cooling stop temperature ± 100 ° C for 10s to 100s after the accelerated cooling stop, Cooling is performed or, if necessary, second accelerated cooling with a cooling rate of 3 to 100 ° C./s and a stop temperature of 500 ° C. or less is performed. In addition, for the sake of distinction,
Accelerated cooling performed subsequent to the rolling is the first accelerated cooling, and after the accelerated cooling is stopped, the temperature of the steel is set within the accelerated cooling stop temperature ± 100 ° C. for 10 seconds to 100 seconds, and then the second cooling or accelerated cooling is performed. This is called cooling or accelerated cooling.

The steel piece is heated to a temperature not lower than the AC3 transformation point and not higher than 1300 ° C., because if the heating temperature is lower than the AC3 transformation point, untransformed ferrite will remain and the subsequent heat treatment step. However, if the heating temperature is higher than 1300 ° C., the heated austenite grain size becomes excessive, and even after recrystallization by hot working, sufficient uniform fine grains can be obtained. This is because austenite grains cannot be obtained and the toughness may deteriorate. After heating the steel slab to a temperature of AC3 transformation point or higher and 1300 ° C. or lower, hot rolling is performed to obtain a desired plate thickness, and at the same time,
Substantially refine austenite and introduce dislocation into austenite, which are necessary for refinement of transformation structure. In order to make hot rolling effective for the refinement of the transformation structure, at least 9
The cumulative rolling reduction ratio is 3 in the range of 50 ° C to the Ar3 transformation point.
It is essential to include 0% or more of austenite unrecrystallized zone rolling. If the cumulative rolling reduction is less than 30%, the effect of refining the transformation structure is not sufficient, so in the present invention, the cumulative rolling reduction of the unrecrystallized region rolling is set to 30% or more. Further, in order to ensure that the unrecrystallized region rolling of austenite,
The rolling needs to be performed in the range of 950 ° C. to Ar3 transformation point. If the rolling temperature is higher than 0950 ° C, the austenite may be recrystallized and the rolling may not be performed in the non-recrystallized region. Further, if the rolling end temperature is lower than the Ar3 transformation point, transformation occurs before the subsequent accelerated cooling. Occurs, the refinement of the structure becomes insufficient, and high ductility cannot be achieved.
In addition, the transformed ferrite may be processed, and not only the ductility but also the toughness may be significantly deteriorated.

After controlling the austenite state by heating and hot rolling, the first accelerated cooling with a cooling rate of 3 to 100 ° C./s brings the austenite fraction to 20 to 70% from the temperature above the Ar3 transformation point. Go to temperature, then 1
During the period of 0 to 100 s, the temperature of the steel is maintained within the first accelerated cooling stop temperature ± 100 ° C., and then the second cooling is performed. The temperature history after the first accelerated cooling is the refinement of the transformation structure, and a proper two-phase structure of ferrite and a hard phase as a soft phase, which is a prerequisite for a low yield ratio and a high uniform elongation, and a hard phase. Is the most important manufacturing process in achieving the refinement of the steel and consequently the formation of retained austenite in the hard phase.

First, the cooling is divided into two stages, and it is difficult to simultaneously achieve the microstructure refinement and the two-phase microstructure necessary for ensuring the toughness only by continuous accelerated cooling from the austenite region to the ferrite region. It is also difficult to secure sufficient retained austenite in the hard phase, which is important for achieving high uniform elongation. That is, the transformed structure is refined by the first accelerated cooling, and untransformed austenite is brought into a supercooled state. Then, during the first and second cooling, a finer ferrite transformation and / or bainite transformation is generated from the supercooled austenite to determine the final two-phase microstructure state, and further untransformed austenite C Thickens and stabilizes. Then, the final cooling is performed to transform the untransformed austenite into a hard phase containing retained austenite. The final cooling is accelerated cooling if necessary.

The first accelerated cooling following the rolling is Ar3.
From the temperature above the transformation point, the austenite fraction is 20 to 70
The cooling rate is 3 to 100 ° C./s until the temperature reaches%. The reason why the accelerated cooling start temperature is set to the Ar3 transformation point or higher is that if the transformation start is less than Ar3, transformation occurs before accelerated cooling and coarse ferrite is generated, which is not preferable for toughness. The accelerated cooling has an austenite fraction of 20.
It is necessary to stop at a temperature of ~ 70%. This is because if the austenite fraction is less than 20%, the outline of the transformation structure is determined in the first accelerated cooling stage, and the retained austenite that can be formed for the first time in the two-stage cooling becomes insufficient. This is because the strength / uniform elongation balance is not optimized because the ferrite phase is too large. On the other hand, when the austenite fraction exceeds 70%, the concentration of C in the austenite is insufficient. Therefore, it is highly possible that transformation proceeds before the second cooling and coarse ferrite or bainite is generated, and in that case, a hard phase suitable for uniform elongation is not formed and high uniform elongation characteristics are obtained. This is because it cannot be achieved.

The cooling rate of the first accelerated cooling is 3 to 100.
C / s. This is because if the cooling rate is less than 3 ° C./s, coarse ferrite transformation occurs during accelerated cooling, or even if transformation does not occur, supercooling of austenite is insufficient and it is effective for microstructuring. This is not preferable, and on the other hand, if the cooling rate exceeds 100 ° C./s and is excessively high, it is industrially difficult to control the accelerated cooling stop temperature, which is not preferable. The interval from the first stop of accelerated cooling to the start of second cooling or accelerated cooling is 10
~ 100s. During this time, any means such as cooling by cooling or holding by an external heating device may be used, but at least the second accelerated cooling start temperature is the first accelerated cooling stop temperature ± 100.
It is necessary to control the temperature to be within ℃.

During the two times of accelerated cooling, supercooled austenite undergoes a finer ferrite transformation and / or
And a step of causing bainite transformation to determine the final two-phase structure state and further concentrating and stabilizing C in untransformed austenite, which is a very important step in the present invention. The interval from the first stop of accelerated cooling to the start of second accelerated cooling is limited to 10 to 100 s, during this time, the temperature of the steel is the first accelerated cooling stop temperature ±
Even if the temperature is within 100 ° C., if the interval is less than 10 s, the transformation does not proceed and the concentration of C in the untransformed austenite becomes insufficient, while if it exceeds 100 s, the second accelerated cooling starts. This is because the austenite may start to transform before. In the former case, the retention of retained austenite in the hard phase cannot be expected, and as a result, improvement in uniform elongation cannot be expected.In the latter case, coarse ferrite or bainite structure is generated, resulting in toughness and uniform elongation. And both deteriorate.

The interval from the stop of the first accelerated cooling to the start of the second cooling is set to 10 to 100 s, and during that period,
Do not exceed the first accelerated cooling stop temperature ± 100 ° C. However, the temperature may vary as long as it is within the first accelerated cooling stop temperature ± 100 ° C. from the first accelerated cooling stop to the second accelerated cooling. Ie 1
The period from the stop of the second accelerated cooling to the start of the second accelerated cooling may be monotonically maintained, the temperature may be increased or decreased, or the temperature may be changed periodically. However, if the cooling rate is too high when the temperature is low, the concentration of C in the untransformed austenite may be insufficient. Therefore, in order to ensure the enrichment of C in the untransformed austenite, The speed should be limited to 0.5 ° C / s or less. If the temperature of the steel decreases from the first accelerated cooling stop temperature by more than 100 ° C. before starting the second cooling, there is a great risk that bainite transformation will occur before the second cooling. Bainite formed during holding or cooling that is not during accelerated cooling becomes coarse and its hardness is not sufficient, which is not preferable for both uniform elongation and toughness. On the other hand, if the temperature of the steel rises from the first accelerated cooling stop temperature by more than 100 ° C. before starting the second cooling, the austenite that has been supercooled is reheated to a high temperature, and as a result, it becomes coarse at high temperature. Since such a large amount of ferrite will be generated, the toughness will be greatly deteriorated, which is not preferable.

The second cooling is not particularly limited as long as it is a cooling rate at which coarse ferrite or bainite is not formed, and may be about air cooling as long as the plate thickness is 50 mm or less. However, in order to surely form a desired hard phase regardless of the chemical composition, accelerated cooling is more preferable depending on the required strength of steel. In particular, for a thick material having a thickness of more than 50 mm, accelerated cooling is preferable in order to ensure the refinement of the structure. When the second cooling is accelerated cooling, the cooling rate needs to be 3 to 100 ° C./s and the stop temperature needs to be 300 ° C. or lower. This is because if the cooling rate is less than 3 ° C./s, the effect of accelerated cooling is insufficient, and if it exceeds 100 ° C./s, there is a problem in securing the shape of steel. Also,
The reason why the accelerated cooling stop temperature when the second cooling is accelerated cooling is 500 ° C. or less is that if it is more than 500 ° C., coarse ferrite and bainite may remain even if the accelerated cooling is performed. Is.

Further, in the present invention, for the purpose of removing residual stress and adjusting strength of steel, the heating temperature is 500 ° C. or higher, A
Tempering below the C1 transformation point can be performed. If the heating temperature for tempering is less than 500 ° C., the tempering effect is not sufficient, while if it is above the AC1 transformation point, the strength of the hard phase decreases and the strength-uniform elongation balance deteriorates, which is not preferable. . Regarding the holding time of tempering and the cooling conditions, the influence on the material is very small compared to the heating temperature, and it is not necessary to prescribe it in a realistic condition range,
The holding time is preferably 48 hours or less in order to suppress the coarsening of the structure, and the cooling rate is more preferably 0.01 ° C./s or higher in order to prevent deterioration of ductility due to coarsening of cementite. Next, the reasons for limiting the chemical composition in the present invention will be explained. First, C is essential as an effective component for improving the strength of steel and for forming a hard phase. If it is less than 0.01%, it is difficult to secure the strength required for structural steel, and hard phase formation is insufficient. On the other hand, excessive addition of more than 0.2% causes excessive embrittlement of the hard phase and deteriorates both toughness and uniform elongation, so the content was made 0.01 to 0.2%.

Next, Si is an element effective as a deoxidizing element and for ensuring the strength of the base material. Addition of less than 0.01% results in insufficient deoxidation and is disadvantageous in securing strength.
On the other hand, excessive addition of more than 1% forms a coarse oxide, leading to deterioration of ductility and toughness. Therefore, the range of Si is 0.01
-1%. Further, Mn is an element necessary for ensuring the strength and toughness of the base material, and it is necessary to add at least 0.1% or more. However, excessive addition of more than 2% causes deterioration of toughness due to the hard phase and deterioration of toughness and cracking property of the welded portion as well as excessive C content, so the upper limit was made 2%.

Al is an element effective for deoxidation, refinement of the structure through grain refinement of the austenite grain size, etc., and in order to exert the effect, it is necessary to contain 0.001% or more,
If added in excess of 0.1%, a coarse oxide is formed and ductility is extremely deteriorated.
It is necessary to limit the range to 0.1%. N is Al or Ti
It works effectively for refinement of austenite grains in combination with
In order to make the effect clear, it is necessary to contain 0.001% or more. However, if it is added excessively, solute N increases, leading to a decrease in ductility and deterioration of the toughness of the base metal and the weld heat affected zone. The upper limit is 0.01% as an allowable range from the viewpoint of ensuring toughness. P and S are impurity elements, and it is preferable to reduce them as much as possible. P has a remarkable tendency to deteriorate the toughness, and the upper limit is set to 0 as an allowable amount from the viewpoint of ensuring the toughness.
It was set to 02%.

Since S forms MnS and deteriorates the ductility value in particular, S is an element which needs to be reduced especially in the steel sheet which is required to secure the ductility as the subject of the present invention. However, the upper limit of the content is set to 0.01% as the upper limit to which ductility deterioration can be practically tolerated. The above are the basic components of the steel of the present invention, but Ni, Cu, Cr, M are added as necessary for the purpose of increasing the strength of the base metal according to the desired strength level.
One or more of o, W, Ti, V, Nb, Zr, Ta and B can be contained.

First, Ni is a very effective element because it can improve the strength and toughness of the base material at the same time, but it is necessary to contain Ni in an amount of 0.01% or more in order to exert the effect. When the content is high, the strength and toughness are improved, but the effect is saturated even if added in excess of 5%. Therefore, considering the economical efficiency, the upper limit is made 5%. Next, Cu has almost the same effect as Ni, but addition of more than 1.5% causes a problem in hot workability, so the content is limited to 0.01 to 1.5%.

Cr is an element effective for improving the strength of the base material, but 0.01% or more is necessary for producing a clear effect. On the other hand, if added in excess of 2%, toughness deteriorates. Therefore, the range is 0.01 to 2%. Mo is also an element effective in improving the strength of the base material, but 0.01% or more is necessary for producing a clear effect, while if added in excess of 2%, the toughness tends to deteriorate. , 0.01 to 2%. W, like Mo, is an element effective for improving the strength of the base material, but 0.01% or more is necessary for producing a clear effect.
If added over 2%, the toughness tends to deteriorate, so the content is made 0.01 to 2%.

[0026] Ti is an element effective in improving strength and refining austenite grains by forming a carbonitride of Ti, and contributes to improving the toughness of the base metal and the weld heat affected zone. In order to form a carbonitride and exert its effect, 0.003% or more must be added. On the other hand, if it exceeds 0.1%, coarse oxides or carbonitrides are formed to deteriorate toughness and ductility, so the upper limit is made 0.1%. V can increase the strength by precipitation strengthening. To be effective, 0.
It is necessary to add 005% or more. On the other hand, if it exceeds 0.5%, coarse carbonitrides are formed and the toughness and ductility are deteriorated. Therefore, in the present invention, the V content is 0.005 to 0.5.
Limit to the range of%.

Nb also has the same effect as Ti or V in the present invention. 0.0 to be effective
It is necessary to add more than 03%. On the other hand, if it exceeds 0.1%, precipitation embrittlement becomes remarkable, and coarse carbonitride is formed to further deteriorate toughness and ductility. Therefore, in the present invention, the Nb content is 0.003 to 0. Limit to 1% range. Zr is also an element that exerts an effect on precipitation strengthening and grain refinement, but in order to exert the effect, 0.003% or more must be added. On the other hand, if over 0.1% is added, the toughness deteriorates due to coarsening of precipitates, so 0.003%
It is limited to the range of 0.1%.

Similarly, Ta is also effective for precipitation strengthening and grain refining, but 0.005% or more is necessary for exhibiting the effect, and if it exceeds 0.2%, toughness deteriorates conversely. The range is 0.005% to 0.2%. B is 0.000
Addition of a very small amount of 2% or more enhances the hardenability of the steel material and is extremely effective for increasing the strength, but if added in excess, it forms BN, which adversely reduces the hardenability and greatly deteriorates the toughness. Therefore, the upper limit is made 0.005%.

Further, in the present invention, one or more of Mg, Ca, Y, La, and Ce are contained for the purpose of stably improving ductility and weld zone toughness (HAZ toughness). be able to. Both improve the ductility characteristics by finely dispersing oxides and sulfides, and refine the structure of the welding heat affected zone (HAZ) to improve the HAZ toughness, but in order to exert its effect, Mg, Ca
Is 0.0005% or more, Y is 0.005% or more, La,
Ce needs to be contained in an amount of 0.001% or more. on the other hand,
If added excessively, the oxides and sulfides become coarse, which itself becomes the starting point of brittle fracture and conversely deteriorates the HAZ toughness, so the upper limits are 0.01% for Mg and Ca, and Y, L
a and Ce are limited to 0.1%.

The reasons for limiting the individual components in the present invention have been described above. In the present invention, however, the carbon equivalent (Ceq.) Represented by the formula (A) must be limited to 0.3 to 0.6%. There is. This is because if the carbon equivalent is less than 0.3%, it becomes difficult to supercool the austenite and stably secure sufficient retained austenite in the hard phase because the hardenability of the steel is too small. On the other hand, on the other hand, if it exceeds 0.6%, on the contrary, the hardenability becomes excessively large, and it becomes difficult to form a two-phase structure of the soft phase and the hard phase in the manufacturing process of the present invention. .

[0031]

The above is a description of the requirements of the present invention. Furthermore, the effects of the present invention will be shown based on Examples. Steel plates having the chemical compositions shown in Table 1 were used to manufacture steel sheets by the method shown in Table 2. In Table 1, steel slab numbers 1 to 12 satisfy the chemical composition of the present invention, and steel slab number 13
Nos. 17 to 17 do not satisfy the chemical composition of the present invention. Table 3 shows the results of examining the tensile properties of the steel sheets shown in Table 2 and the toughness by the 2 mm V notch Charpy impact test. Steel plate symbols A1 to A15 in Tables 2 and 3 are examples satisfying all the requirements of the present invention, and steel plate numbers B1 to B11 do not satisfy the requirements of the chemical composition or / and the manufacturing method of the present invention. This is a comparative example.

In the production of the steel sheet, water was used for the first and second accelerated cooling. The temperature control between the first accelerated cooling and the second (accelerated) cooling is
In the holding or heating process, a high-frequency induction heating device was used, and in the case of cooling, it was left to cool to a predetermined temperature as it was or inserted into a heat treatment furnace set to an appropriate temperature and then cooled in the furnace. The mechanical properties were measured by taking a test piece from the center of the plate thickness at right angles to the rolling direction. The tensile test piece was a round bar test piece having a parallel portion of 6 mmφ × 24 mm, and the Charpy test piece was a standard test piece having a test piece thickness of 10 mm. The tensile test was carried out at room temperature, and the Charpy test was carried out at various temperatures to determine the fracture surface transition temperature (vTrs).

Regarding the mechanical properties shown in Table 3, the results of comparing the relationship between the tensile strength and the uniform elongation between the inventive example and the comparative example are shown in FIG. The tensile strength / uniform elongation balance of the steel sheet manufactured according to the present invention is extremely good as compared with the tensile strength / uniform elongation balance of the comparative example by the conventional manufacturing method or the composition not satisfying the requirements of the present invention and the manufacturing method. Is clear. Note that among the comparative examples, there is a steel sheet having a tensile strength / uniform elongation balance close to that of the present invention example, but the toughness of the steel sheet is extremely poor. It is clear that all good high strength steels can be produced.

The reason why each comparative example is inferior to the present invention will be described in more detail below. First,
Steel plate numbers B1 to B5 are examples in which the chemical composition does not satisfy the present invention. Steel plate No. B1 does not satisfy the present invention and has an excessively large amount of C, so the manufacturing method satisfies the present invention, but the hard phase is significantly embrittled, and the toughness and uniform elongation are both the steels of the present invention. Markedly inferior to.

Steel plate No. B2 has an excessively large amount of Mn, so that the toughness of the hard phase is greatly deteriorated, and both the toughness and the uniform elongation are markedly inferior to those of the steels of the present invention. Steel plate number B3
Has an excessively large amount of S that adversely affects the ductility, so the toughness is slightly inferior, but the uniform elongation is significantly deteriorated. Steel plate No. B4 has a carbon equivalent that is too small and the hardenability of the steel is not sufficient, so that a hard bainite structure containing retained austenite and having a high strength is not formed, and a coarse bainite structure is formed. , Uniform elongation is low for strength, and toughness is poor. Steel plate number B5, on the other hand, has an excessively large carbon equivalent, so that the formation of soft ferrite is insufficient, and a clear two-phase organization of the soft phase and the hard phase is not achieved. Therefore, toughness and uniform elongation are achieved. Both are greatly deteriorated.

Next, steel sheet numbers B6 to B11 are examples in which the requirements regarding the manufacturing method do not satisfy the present invention. Steel plate number B6 is a normal thermo-mechanical treatment step for accelerated cooling from the austenite region to the completion of transformation, and therefore, a two-phase microstructure consisting of a soft phase and a hard phase, which is necessary for improving uniform elongation, is present in the hard phase. The formation of residual austenite of 1 is not achieved, and the uniform elongation is clearly inferior to the steels of the present invention having the same strength level. Steel plate No. B7 had a first accelerated cooling start temperature below the Ar3 transformation point, so coarse ferrite was generated before accelerated cooling, and because it was not in two-stage accelerated cooling, The refinement and formation of retained austenite are not sufficient, the toughness is significantly deteriorated, and the uniform elongation is not improved as much as the present invention.

Steel plate No. B8 had an excessively large austenite fraction in the first accelerated cooling stop stage, so that coarse ferrite was produced before the second accelerated cooling, and the C concentration in austenite was increased. Since the hardness is insufficient, the hardness of the hard phase is insufficient, and the formation of retained austenite in the hard phase is insufficient. Therefore, both toughness and uniform elongation are inferior. Steel plate number B9 is the first accelerated cooling and 2
This is an example in which the temperature history during the accelerated cooling of the third time is within the scope of the present invention. That is, before starting the second accelerated cooling, the steel plate temperature is 100 ° C. from the first accelerated cooling stop temperature.
Since the temperature has reached an extremely high temperature, coarse ferrite is generated and the toughness is greatly deteriorated. Since the structure is coarse, uniform elongation is insufficient.

Steel plate number B10 is also an example in which the temperature history between the first accelerated cooling and the second accelerated cooling is within the scope of the present invention. That is, this comparative example is a process of monotonically cooling after the first accelerated cooling and then entering the second accelerated cooling, but the temperature decrease until the second accelerated cooling enters the range of the present invention. Is too large, and the interval between the first accelerated cooling and the second accelerated cooling is too large. Therefore, coarse bainite, which is not preferable for toughness and uniform elongation, is partially formed before the second accelerated cooling, and both toughness and uniform elongation are poor. Steel plate number B11 is an example in which the temperature between the first accelerated cooling and the second accelerated cooling is held at a constant temperature. However, since the holding time is too long, a partial transformation occurs during the holding and coarse ferrite, A hard phase is formed, and both toughness and uniform elongation are insufficient as compared with the present invention. Also from the above examples, according to the present invention, while maintaining good low temperature toughness, uniform elongation is excellent in comparison with the same tensile strength, production of high toughness / high ductility steel has good productivity. It is clear that this is possible in the thermo-mechanical treatment step.

[Table 1]

[Table 2]

[Table 3]

[0039]

EFFECTS OF THE INVENTION According to the present invention, high toughness with high safety, having sufficient strength as a welded structural steel, excellent ductility characteristics such as uniform elongation, and low temperature toughness. -High-ductility high-strength steel can be manufactured in a highly productive thermomechanical process without depending on the addition of a large amount of expensive alloy elements, and the industrial effect is extremely remarkable.

[Brief description of drawings]

FIG. 1 is a schematic view showing an outline of a manufacturing process of the present invention.

FIG. 2 is a diagram comparing the example of the present invention with the comparative example in terms of the relationship between the tensile strength and the uniform elongation regarding the mechanical properties shown in the examples.

   ─────────────────────────────────────────────────── ─── Continued front page    (72) Inventor Hiroyuki Shirahata             No. 1 Nishinosu, Oita-shi, Nippon Steel Corporation             Company Oita Works F-term (reference) 4K032 AA01 AA02 AA04 AA05 AA08                       AA11 AA12 AA14 AA15 AA16                       AA21 AA22 AA23 AA24 AA27                       AA29 AA31 AA33 AA35 AA36                       AA37 AA39 AA40 BA01 BA02                       BA03 CA02 CA03 CB01 CB02                       CC03 CC04 CD03

Claims (5)

[Claims] 1. In mass%, C: 0.01 to 0.2%, Si: 0.01 to 1%, Mn: 0.1 to 2%, Al: 0.001 to 0.1%, N : 0.001 to 0.01%, as impurities, P: 0.02% or less, S: 0.01% or less, and the carbon equivalent (Ceq.) Represented by the following formula (A) is 0.3 ~
A steel slab containing 0.6% of balance Fe and unavoidable impurities at 0.6% is heated to a temperature of AC3 transformation point or higher and 1300 ° C. or lower,
At least a start temperature of 950 ° C or lower and an end temperature of Ar3
After performing hot rolling including unrecrystallized region rolling of austenite having a cumulative reduction of 30% or more at the transformation point or higher, the cold speed is 3%.
Accelerated cooling up to 100 ° C / s is performed from a temperature of Ar3 transformation point or higher to a temperature at which the austenite fraction is 20 to 70%, and after accelerated cooling is stopped, temperature is raised, held, and the cooling rate is 0.5 ° C / s or less One or a combination of two or more types of cooling is performed, and after the accelerated cooling stop, the temperature of the steel is maintained within the accelerated cooling stop temperature ± 100 ° C. for 10 s to 100 s, followed by cooling. Of high toughness, high ductility and high tensile strength steel. Ceq. = C% + Mn% / 6 + (Cu% + Ni%) / 15+ (V% + Mo% + Cr%) / 5 ... (A) 2. The temperature of the steel is accelerated cooling stop temperature ± 10.
After maintaining the temperature within 0 ° C, the cooling rate was further increased to 3 to 100 ° C /
The high toughness according to claim 1, wherein the second accelerated cooling is performed at s and the stop temperature is 500 ° C or less.
Manufacturing method of high ductility and high strength steel. 3. A heating temperature of 500 ° C. or higher, AC
A tempering process at a temperature not higher than one transformation point is applied.
Alternatively, the method for producing a high-toughness / high-ductility high-strength steel according to any one of 2) to 3). 4. Further, in mass%, Ni: 0.01 to 5%, Cu: 0.01 to 1.5%, Cr: 0.01 to 2%, Mo: 0.01 to 2%, W : 0.01-2%, Ti: 0.003-0.1%, V: 0.005-0.5%, Nb: 0.003-0.1%, Zr: 0.003-0.1% %, Ta: 0.005-0.2%, B: 0.0002-0.005%, 1 type, or 2 or more types are contained, In any one of Claim 1 thru | or 3 characterized by the above-mentioned. Of high toughness, high ductility and high strength steel. 5. Further, in mass%, Mg: 0.0005-0.01%, Ca: 0.0005-0.01%, Y: 0.005-0.1%, La: 0.001- 0.1%, Ce: 0.001-0.1%, 1 type (s) or 2 or more types are contained, The high toughness and high ductility of Claim 1 characterized by the above-mentioned. Method of manufacturing tensile steel.