JP3499085B2 - Low Yield Ratio High Tensile Steel for Construction Excellent in Fracture Resistance and Manufacturing Method Thereof - Google Patents

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a steel material for a strength member used for a structure which is subjected to a large and repeated plastic strain due to a large earthquake or the like during use and a method for manufacturing the same. For example, the steel material produced by this method can be generally used for welded steel structures such as offshore structures, pressure vessels, shipbuilding, bridges, buildings, line pipes, etc. It is useful as a steel material for structures requiring construction, such as bridges. Further, the form of the steel material is not particularly limited, but it is used as a structural member,
It is particularly useful for steel plates requiring low temperature toughness, particularly thick plates, steel pipe materials, or shaped steel.

[0002]

2. Description of the Related Art In recent years, structures have tended to increase in size as seen in higher-rise buildings, larger spans of bridges, etc. Ensuring performance to prevent the collapse of objects has become an important issue. In particular, from the experience of the Great Hanshin Earthquake, if the safety of a structure is to be ensured by the performance of steel without special consideration in design and construction, it is necessary to have a steel with high safety in terms of both ductile fracture and brittle fracture. It is being recognized.

Recently, use of a low yield ratio steel (low YR steel) or a high uniform elongation steel in consideration of ductile fracture performance has been investigated for high-rise building steel materials. It has been recognized that the low yield ratio characteristics are excellent in the ability to absorb energy due to earthquakes, typhoons, etc., and are also useful in suppressing the local collapse of structures.

If the collapse of a structure due to an earthquake is caused only by ductile fracture of the material, the use of such a steel material leads to an improvement in the safety of the structure. However, in a huge earthquake such as the Great Hanshin Earthquake, steel materials do not always have a ductile fracture and finally collapse, and a ductile fracture is followed by a brittle fracture, and the brittle crack propagates throughout. Various post-earthquake studies have shown that it may cause eventual structural collapse.

Further, it is appropriate to think that the deformation due to an earthquake is not simple, and that particularly in the case of a huge earthquake, the huge and continuous vibrations repeatedly apply a large level of force that causes plastic deformation of the steel material. is there.

From the above viewpoints, in order to prevent the structure from collapsing due to a huge earthquake or huge typhoon such as once every several decades to hundreds of years, in addition to the low yield ratio characteristics excellent in energy absorption capacity, Of the properties additionally required for the steel material, it is important to add the following properties (1) to (3).

Due to the improved ductility characteristics, the ductile fracture capacity capable of absorbing the energy of earthquake is excellent, and the ductile crack is generated and the propagation resistance is large. Little deterioration in toughness due to repeated plastic deformation. Even if a brittle crack occurs once, it stops halfway and does not lead to the destruction or collapse of the entire members and structures. Conventionally, various techniques have been developed for securing the above characteristics, and techniques for partially improving the individual characteristics have been developed.

Many means of lowering the yield ratio for improving the energy absorption capacity have been proposed. For example, a method by adjusting the chemical composition such as an increase in the amount of C, a method of coarsening crystal grains, a method of performing an intermediate heat treatment of heating in a ferrite (α) + austenite (γ) two-phase region between quenching and tempering heat treatments. As represented by (hereinafter, abbreviated to QLT treatment), there is a method of mixing α as a soft phase and bainite or martensite as a hard phase.

[0009] For example, as a manufacturing method for obtaining a mixed structure of a soft phase and a hard phase, Japanese Patent Laid-Open No. 23817/1993 discloses that a steel sheet is reheat-quenched and then Ac ₁ transformation point and Ac ₃
A method of reheating between transformation points to obtain two phases of γ and α and then air cooling is disclosed, and further, it is disclosed in JP-A-4-31482.
Similarly, Japanese Patent Publication No. 4 discloses a method of quenching after reheating to a two-phase region.

Further, as a method for producing on-line without performing reheating treatment, for example, JP-A-63-286517
The publication discloses a method in which after hot rolling from the γ region to the two-phase region, air cooling is performed to a temperature 20 to 100 ° C. lower than the Ar ₃ transformation point to generate the α phase, and then the material is rapidly cooled.

It has been conventionally known that the inclusion of Ni is effective for stopping brittle cracks. Further, recently, a technique has been developed, as disclosed in Japanese Patent Application Laid-Open No. 4-141517, for improving the propagation stopping property of a brittle crack without increasing the amount of Ni by imparting an ultrafine grain structure to the surface layer portion. ing.

[0012]

However, before the experience of the Great Hanshin Earthquake, the above-mentioned characteristics (1), which were not conceived as necessary steel material characteristics, in particular, are not always sufficiently recognized, and therefore, The characteristics of ~ are satisfied at the same time, and even if a huge earthquake or huge typhoon such as once every several decades to several hundred years is encountered without any special consideration in design and construction, the structure will not collapse. It cannot be said at this stage that steel materials with excellent fracture resistance and their manufacturing technology have been established.

From the viewpoint of ensuring safety as a structure,
Naturally, the first consideration is to control the occurrence of brittle fracture. However, from the viewpoint of ensuring safety in the event of a huge earthquake, which is rarely seen, it is necessary to suppress the initiation and propagation of ductile cracks that facilitate the initiation of brittle fracture, and to prevent crack tip after ductile crack growth. It is a new problem imposed on the steel sheet that no large deterioration of toughness in the plastic region and no large deterioration of toughness after repeated plastic deformation occur.

However, the structure has a welded portion with deteriorated characteristics, and it is impossible to completely eliminate the existence of a defect which becomes a starting point of brittle fracture. In that case, it is extremely difficult to completely suppress the occurrence of brittle fracture itself, and it is also extremely economically disadvantageous. Therefore, it is necessary to allow the brittle fracture in the unlikely event and to make the propagation stopping property of the brittle crack capable of preventing the propagation of the crack compatible with the ductile fracture property and the toughness after plastic deformation.

A technique for forming a superfine grain structure in the surface layer to improve the propagation stopping property of a brittle crack is disclosed in Japanese Patent Application Laid-Open No. 4-141.
As disclosed in Japanese Laid-Open Patent Publication No. 517, etc., when a large plastic deformation due to a large earthquake, etc., which is the object of the present invention, toughness and ductility are obtained only by forming an ultrafine grain structure in the surface layer portion. It is difficult to secure resistance to fracture, such as brittle crack propagation arresting properties.

Further, Japanese Laid-Open Patent Publication No. 4-141517 does not disclose a technique for imparting a low yield ratio, which has a great effect on improving the seismic resistance. In order to form a superfine grain structure in the surface layer with a low yield ratio. New technology is required.

[0017]

The inventors of the present invention need to reduce P and S as impurities and O (oxygen) as much as possible in order to improve the characteristics of ductile fracture initiation and crack propagation arrest. It is important to reduce the amount of solute N in order to maintain its properties even after repeated plastic deformation, and to reduce the toughness deterioration due to plastic strain such as repeated plastic deformation, reduce the amount of solute N in steel. I thought I needed to.

To this end, Al, Ti, Zr, Nb, Ta, V, B which has the effect of forming a nitride to fix N.
It has been found through detailed experimental analysis that it is important to properly contain steel, and it has both ductile fracture characteristics, brittle fracture initiation characteristics (toughness) and propagation arrest characteristics of brittle cracks regardless of the steel type and composition range. It has been found that the most preferable means for improving the propagation stopping property of brittle cracks is to impart an ultrafine grain structure to the surface layer.

In particular, the average ferrite grain size is 3 μm.
With an ultrafine grain structure of m or less, even when subjected to severe plastic deformation such that the plastic strain exceeds 10%,
It was clarified that the deterioration of the Charpy impact properties and the propagation arrest properties of brittle cracks was significantly smaller than that of ordinary structures.

Further, in order to secure the brittle crack propagation arresting property after plastic deformation, it is necessary to suppress not only the surface layer portion but also the internal toughness due to the plastic deformation to some extent. It was also clarified experimentally that the grain size needs to be reduced.

Further, it has been experimentally confirmed that the reduction of the amount of N, the fixing of N by the nitride forming element, and the addition of the ultrafine grain layer to the surface layer portion are effective for improving the toughness and ductility of the welded joint. .

By comprehensively analyzing the above findings, the present invention improves ductility by limiting chemical components and deteriorates toughness due to plastic deformation, and further forms a superfine grain structure in the surface layer to form brittleness. It has been found that in a steel having improved crack propagation arresting properties, a means for dispersing an appropriate amount of martensite phase in the structure is the most preferable means for obtaining the low yield ratio property.

That is, although various means for obtaining a low yield ratio characteristic are conceivable, when the superfine grain structure is present in the surface layer, deterioration of toughness is suppressed even if a brittle hard phase such as martensite is dispersed. Therefore, it is most suitable to use the low yield ratio due to the dispersion of martensite, in which the steel composition is relatively limited. The internal toughness deterioration is inevitable by the method of increasing the grain size of the inside of the steel excluding the surface layer to lower the yield ratio of the entire steel material.

As described above, according to the present invention, it is found that the dispersion of martensite is the most preferable means for lowering the yield ratio in the steel material having the surface ultrafine grain layer structure, and the surface layer The inventors have established and invented a manufacturing method capable of simultaneously achieving ultrafine grain structure and dispersion of martensite, and the gist thereof is as follows.

(1) C: 0.01 to 0.1 by mass %
5%, Si: 0.01 to 1.0%, Mn: 0.1 to 2.
0%, Al: 0.003-0.1%, N: 0.001-
0.006% and N (%)-Al (%) /
When 3.0 ≦ 0, the contents of P, S and O as impurities are
P: 0.01% or less, S: 0.01% or less, O: 0.0
A steel material containing less than 06% and balance iron and unavoidable impurities, having an average crystal grain size of 30 μm or less at the center of the plate thickness, and having a martensite ratio in the steel material volume of 1
0 to 60%, and at least two of the outer surfaces constituting the steel material have an average ferrite grain size within a range of 10 to 33% of the total thickness from the surface layer to 3 μm.
A low-yield ratio, high-strength steel material for construction with excellent fracture resistance characterized by the following ultrafine grain structure.

(2) Ti: 0.003% by mass %
0.020%, Zr: 0.003 to 0.10%, Nb:
0.002-0.050%, Ta: 0.005-0.2
0%, V: 0.005 to 0.20%, B: 0.0002
To 0.003%, one or more of
(%)-Al (%) / 3.0-Ti (%) / 3.4-Z
r (%) / 6.5-Nb (%) / 13.2-Ta (%)
/25.8-V(%)/10.9-B(%)/2.0≤
The low yield ratio high tensile steel material for construction having excellent fracture resistance as described in (1) above, which is 0.

(3) In mass %, Cr: 0.01 to 2.
0%, Mo: 0.01 to 2.0%, Ni: 0.01 to
4.0%, Cu: 0.01 to 2.0%, W: 0.01 to
2.0% of 1 type (s) or 2 or more types, The low yield ratio high tensile steel material for construction excellent in the fracture resistance performance of said (1) or (2) characterized by the above-mentioned.

(4) Mg: 0.0005 to 5% by mass
0.01%, Ca: 0.0005-0.01%, RE
M: 0.005 to 0.10% of 1 type or 2 types or more are contained, The construction excellent in the fracture resistance performance of any one of said (1)-(3) characterized by the above-mentioned. Low yield ratio, high strength steel.

(5) A steel slab having the composition described in any of (1) to (4) above is heated to a temperature of Ac ₃ transformation point or higher and 1250 ° C. or lower, and accumulated in an austenite region of 950 ° C. or lower. After performing a rough rolling with a rolling reduction of 10 to 50%, at least two surface layer regions of the outer surface corresponding to 10 to 33% of the thickness of the steel slab at that stage are changed from a temperature of Ar ₃ transformation point or higher to 2 to In the process of starting the cooling at a cooling rate of 40 ° C./s and stopping the cooling at the Ar ₃ transformation point or lower to reheat the heat at least once, the reheat after the last cooling from the start of the cooling is Finish rolling with a cumulative reduction of 20 to 90% is completed until the end, and the surface layer area of the steel material after the rolling is completed is (Ac ₁ transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.).
After reheating to the range of 0, the steel material after the reheating is further reduced to 0.
2 to 2 ° C. / s in the cooling rate after cooling to a range of (the cooling rate in transformation start temperature _{(Ar 3) -50 ℃) ~500} ℃, 20~300 ℃ at a cooling rate of 5 to 40 ° C. / s To produce a steel material according to any one of (1) to (4) above, and a method for producing a high-strength steel material with a low yield ratio for construction having excellent fracture resistance.

(6) A steel slab having the composition described in any of (1) to (4) above is heated to a temperature of Ac ₃ transformation point or higher and 1250 ° C. or lower, and accumulated in an austenite region of 950 ° C. or lower. After performing a rough rolling with a rolling reduction of 10 to 50%, at least two surface layer regions of the outer surface corresponding to 10 to 33% of the thickness of the steel slab at that stage are changed from a temperature of Ar ₃ transformation point or higher to 2 to In the process of starting the cooling at a cooling rate of 40 ° C./s and stopping the cooling at the Ar ₃ transformation point or lower to reheat the heat at least once, the reheat after the last cooling from the start of the cooling is Finish rolling with a cumulative reduction of 20 to 90% is completed until the end, and the surface layer area of the steel material after the rolling is completed is (Ac ₁ transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.).
After recuperating within the range, the steel after the recuperation is allowed to cool, or the steel after the recuperation is cooled to 20 to 650 ° C at a cooling rate of 5 to 40 ° C / s, and then 0 1-5
(Ac ₁ transformation point + 10 ° C) ~ (Ac
₃ transformation point -30 ℃) in the range of 1 to
After being held for 60 s, a two-phase region heat treatment of cooling at 0.5 to 50 ° C./s is performed to produce the steel material according to any one of (1) to (4) above, which is characterized by fracture resistance. A method of manufacturing a high-strength, low-yield-strength steel material with excellent performance.

(7) A method for producing a low-yield ratio, high-strength steel material for construction having excellent fracture resistance as described in (5) or (6) above, which comprises performing tempering at 450 to 650 ° C. The high-strength steel material here means not only high-tensile steel plates (thick plates) but also steel materials including shaped steel and pipe materials.

[0032]

BEST MODE FOR CARRYING OUT THE INVENTION The requirements for chemical components in the present invention are that the amounts of P, S, and O as impurities for improving ductility after plastic deformation are limited, and that the toughness is deteriorated due to huge and repeated plastic deformation. The limitation is on the chemical composition for the fixation of N for the suppression of N. Hereinafter, these requirements will be described in detail first.

In order to improve the plastic deformability, suppress the initiation and growth of ductile cracks, it is necessary to increase the ductility of the ferrite matrix of the steel. For that purpose, the solid solution P, C and N should be reduced. Is preferred. As for C, the solid solubility limit in ferrite is small and precipitates are easily formed, so that the adverse effect on the ductility characteristics of practical steel can be ignored. Further, since C is an essential element for securing strength, it is not preferable to completely remove it.

N is effective for refining the heated austenite grain size by the nitride, and its inclusion as an impurity is unavoidable, but unlike C, practical steel also forms a solid solution in the ferrite matrix in a certain amount and becomes ductile. It adversely affects the characteristics. Furthermore, when the steel undergoes plastic deformation after plastic deformation or ductile crack growth in the presence of solute N, the toughness becomes remarkable due to the interaction with the dislocations caused by the plastic deformation and the fine precipitation on the dislocation lines. Since it deteriorates, the solid solution N should be reduced to the limit.

For that purpose, it is necessary to fix N by a nitride forming element. As the nitride-forming element, Al is preferable because it has the least effect on other properties, and 10% is added to the superfine grain structure of the surface layer, which is the most important for the propagation stopping property of brittle cracks.
When the Al amount necessary for the deterioration of the toughness due to the addition of the plastic strain to be 20 ° C. or less due to the increase of the fracture transition temperature in the Charpy test is experimentally obtained, the relationship as shown in the equation (1) is obtained. Was obtained. Therefore, in the present invention, the amounts of N and Al are limited to the relationship of the formula (1) on the assumption that the contents of Al and N are within the limited range for the reason described later.

[0036]         N (%)-Al (%) / 3.0 ≦ 0 …………………… (1)

In addition to Al, one or more of Ti, Zr, Nb, Ta, V and B can be selectively used as the nitride forming element. In that case, Al,
The contents of Ti, Zr, Nb, Ta, V, and B are related to the amount of N necessary for suppressing the toughness deterioration due to plastic deformation under the same criteria as in the formula (1) ((( Since it is necessary to adjust the content so that equation 2) is established,
In relation to the amount of N, Ti, Zr, Nb, Ta, V and B are limited so that the relation of equation (2) is established. N (%)-Al (%) / 3.0-Ti (%) / 3.4-Zr (%) / 6.5-Nb (%) / 13.2-Ta (%) / 25.8 -V (%) / 10.9-B (% ) /2.0≦0 …………… (2)

It is also important to limit the amount of P as an impurity. That is, since P deteriorates the ductility of the ferrite matrix, it is necessary to limit its content in order to improve the plastic deformability, the occurrence of ductile cracks, and the growth characteristics. The smaller the amount of P is, the more preferable it is. However, reducing the amount of P puts a load on the refining process, resulting in a decrease in productivity and an increase in cost. Based on this, the content is made 0.01% or less.

That is, the degree of deterioration of the ductility characteristics with the increase of the amount of P becomes remarkable when it exceeds 0.01%. When the amount of P is 0.01% or less, the adverse effect of P is small. Therefore, in the present invention, P as an impurity is used.
The amount is limited to 0.01% or less. However, when it is used in a structure where local plastic deformation in the segregation part or deterioration of ductile fracture characteristics has an effect, if the problem of refining is ignored, the P content is 0.007% or less. It is more preferable to limit to

Since S forms MnS, it deteriorates the ductile fracture characteristics. In particular, it deteriorates the propagation characteristics of ductile cracks. Since the ductile fracture initiation characteristics deteriorate under the condition of a large amount of solid solution P and N, deterioration of the ductile crack propagation characteristics due to S greatly affects the plastic deformability and ductile fracture characteristics of the entire steel material. It becomes necessary to reduce S to 10 ppm or less.

However, if the amount of P and N is reduced and the solid solution N is fixed by a nitride forming element as in the present invention,
Since the allowable amount of S widens because the resistance until the occurrence of ductile fracture becomes large, the present invention limits S as an impurity to 0.01% or less based on the experimental results.

Further, it is preferable to reduce O as much as possible in order to form inclusions harmful to ductility. However, like S, ductility of the base material is secured to some extent if the solid solution P and N are reduced. Therefore, the allowable amount is higher than that in the case where the solid solution P and N are not reduced, and it is 0.006 based on the experimental result.
It is necessary to limit it to% or less.

The above is the component limiting range for improving the plastic deformability and the ductile fracture characteristics, which are the requirements of the present invention, and for suppressing the toughness deterioration after the plastic deformation, but it is another important requirement of the present invention. In order to improve the propagation stopping property of brittle cracks, the average ferrite grain size is 3 μm in the surface layer portion of at least two surfaces of the steel material on the premise that the other components will be described in addition to the above-mentioned components. It is necessary to allow the following ultrafine grain structure to exist from the surface layer to a thickness of 10 to 33% of the plate thickness.

By forming an ultrafine grain structure in the surface layer portion, a shear lip which is a ductile fracture is formed in the surface layer portion during the development of the brittle crack, and the brittle crack propagation stopping property is improved. This method is advantageous in that the brittle crack propagation arresting property can be improved without adding or adjusting alloy components.

In order to reliably generate a shear lip by resisting a brittle crack propagating at a high speed, it is necessary to remarkably improve the occurrence of brittle fracture in the surface layer portion and the propagation stoppage property to the required toughness of the steel sheet. Therefore, for that purpose, it is an essential condition that the ferrite grain size of the surface layer portion is remarkably reduced.

Further, in the ultrafine grain structure, deterioration of toughness and brittle crack propagation stopping characteristics due to plastic deformation is very small,
In the steel material that is subject to plastic deformation due to a large earthquake, etc., which is the subject of the present invention, and needs to have toughness and brittle crack propagation stopping properties to the extent that safety can be ensured even after plastic deformation , Is the most advantageous means for improving brittle crack propagation arrest properties.

The ferrite grain size in the surface layer portion is preferably finer as a matter of course. However, in the present invention, the average ferrite grain diameter in the surface layer portion is defined as a range in which the formation of the shear lip is sure and the manufacturing process is not overloaded. The diameter is limited to 3 μm or less.

The ferrite grain structure of the surface layer is preferably a grain size with less variation in crystal grain size.
Even if there are coarse particles having a size of more than twice the average particle size, if the existence ratio is within 10% with respect to the entire surface layer portion, the brittle fracture characteristics of the surface layer portion are not substantially adversely affected. ,
Permissible.

Even if brittle fracture occurs from a defective portion or a welded portion and propagates, the surface layer portion surely undergoes ductile fracture to form a shear lip, and the limitation of the ferrite grain size is an essential condition. However, in order to improve the propagation stopping property of brittle cracks, the thickness of the surface ultrafine grain layer is also an important requirement.

That is, in order to stop the brittle crack of the normal structure inside the steel sheet, it is necessary to absorb the propagation energy at the shear lip portion, but if the thickness of the shear lip is insufficient, the shear lip may be formed. However, the brittle crack may not stop.

The shear lip needs to have a certain thickness in order to surely stop the propagation of the brittle crack. Naturally, the thicker the shear lip, the greater the effect of stopping the cracks, but if you try to secure the thickness of the ultrafine grain layer more than necessary,
An excessive load is applied to the manufacturing process, and depending on the manufacturing conditions, the ductility of the base material, the shape of the steel sheet, the surface quality, etc. may be deteriorated.

In the present invention, the thickness of the superfine grain structure of the surface layer having an average ferrite grain size of 3 μm or less is defined as the range not causing these problems, the lower limit is from the surface layer to 10% of the plate thickness, and the upper limit is the surface layer. To 33% of the extra thickness.

The surface ultrafine grain layer is preferably applied to all surfaces of the steel material, but if the above conditions are satisfied, brittle cracks can be formed by applying the ultrafine grain layers to at least two surfaces. Effective for stopping.

If the surface layer portion has an ultra-fine grain structure, the resistance to the propagation of brittle cracks in the surface layer portion even after plastic deformation does not extremely deteriorate the brittle crack propagation stopping property, but It does not mean that the properties do not contribute at all, and it is also necessary to pay attention to ensuring internal toughness.

Although the internal toughness has an adverse effect of plastic deformation, the grain refining is an effective means for reducing the adverse effect and securing the toughness level. Of course, the finer the internal particle size is, the more advantageous it is, but when the internal structure is miniaturized, it becomes difficult to miniaturize the surface layer portion, or the load on the manufacturing process becomes large. Based on the survey results, the grain size at the center of the plate thickness is limited to 30 μm or less based on the level required to maintain a sufficient level of the brittle crack propagation arresting property of the entire steel material before and after plastic deformation. To do.

The entire grain size other than the superfine grain layer on the surface layer is 30 μm
It is preferable that the following conditions are satisfied. However, when the steel material is broken, the stress condition is the most severe in the plate thickness center part, and in general, the grain size in the plate thickness center part is the coarsest. It defines the crystal grain size at the center of thickness.

The crystal grain size here is a so-called "effective grain size" which can be used as an index showing resistance to fracture.
Indicates. Therefore, the ferrite main structure corresponds to almost the ferrite crystal grain size, and the bainite or martensite main structure corresponds to the bainite packet size and the martensite packet size, respectively.

It should be noted that reduction of the amount of N, fixation of N by a nitride forming element, and addition of an ultrafine grain layer to the surface layer portion, which are effective for securing the toughness after plastic deformation, brittle crack propagation arresting property and ductility, It was experimentally confirmed that it is effective to have an additional effect of improving the toughness and ductility of the welded joint. That is, the solute N has a large adverse effect on the toughness of the heat-affected zone in high heat input welding, but if the amount of N and the amount of solute N are strictly controlled as in the present invention, the amount of solute N will change. The adverse effect on the toughness of the heat affected zone is minimized.

The ultra-fine grain layer formed on the surface layer completely disappears in a region of the heat-affected zone where it is reheated to 1300 ° C. or higher in the vicinity of the weld bead, but is heated to a lower temperature. Although the ultra-fine grain structure disappears in the heat-affected zone, the effect of the ultra-fine grain structure before transformation remains and has the effect of refining the structure of the weld heat-affected zone. Is also effective against.

As described above, it is important to limit the chemical composition in order to improve the ductile fracture characteristics, to improve the propagation stopping characteristics of brittle cracks, and to make the surface layer of the steel material finer, but to secure the fracture resistance performance. As a precondition, the steel must have a low yield ratio.

In the steel in which the propagation arresting property of brittle cracks is improved by forming an ultrafine grain structure in the surface layer, a means for obtaining the low yield ratio property is to use an appropriate amount of martensite phase in the structure. Most preferred is a means to disperse in the tissue.
That is, various means for obtaining a low yield ratio property are conceivable, but when an ultrafine grain structure exists in the surface layer, deterioration of toughness is suppressed even if a brittle hard phase such as martensite is dispersed. It is most suitable when using a low yield ratio due to the dispersion of martensite, which has relatively few restrictions on the steel composition.

Other means for lowering the yield ratio, such as C, Cr,
The increase of the second phase due to the addition of Mo or the like causes an increase in the alloy cost and may have a bad influence on the weldability and the like, and the crystal grain size of the inside excluding the surface layer portion is coarsened to make the steel as a whole. Degradation of internal toughness is inevitable with the method of achieving a low yield ratio. The present invention finds that the dispersion of martensite is the most preferable means for lowering the yield ratio in the steel material having the surface ultrafine grain structure, and the dispersion of the surface ultrafine grain structure and martensite and We have established a manufacturing method that can achieve

The structural requirement for lowering the yield ratio without deteriorating the ductility characteristics is that the ratio of the martensite phase, which is the hard phase, to the steel material volume is 10 to 60%. That is, a means for increasing the tensile strength and lowering the yield ratio (yield stress / tensile strength) by dispersing in the mother phase a second phase having sufficiently higher strength than the mother phase for lowering the yield ratio. Is the most effective.

In the present invention, based on the experimental results, the martensite phase is the most preferable as the hard phase, and the ratio thereof is in the range of 10 to 60% on the average in the volume of the steel sheet, which indicates a low yield ratio and other materials. It has been found that it is most preferable in terms of compatibility with characteristics. Martensite phase ratio is 10
If it is less than%, the effect of improving the tensile strength by the hard phase cannot be obtained, so that a low yield ratio cannot be achieved.

On the other hand, when the ratio of the martensite phase is more than 60%, the concentration of C in the martensite is not sufficient, so that the hardness of martensite decreases and the difference from the hardness of the matrix phase decreases. Therefore, the effect of the martensite phase, which is a hard phase, on the yield stress begins to occur, and the yield ratio increases because the yield stress increases and the tensile strength decreases.

On the other hand, if the proportion of the martensite phase exceeds 60%, coarsening of martensite occurs and ductility and toughness deteriorate, which is not preferable. It should be noted that the martensite phase here contains MA containing a part of retained austenite phase.
It also contains a phase (Martensite-Austenite Constituent).

As a means for obtaining a structure morphology containing a part of a martensite phase, as disclosed in JP-A-53-23817 and the like, austenite ( γ) after re-precipitating the phase,
A typical method is to transform the γ phase into a martensite phase during cooling by cooling by cooling.

However, in the present invention, it is necessary to have a low yield ratio and, at the same time, to have an ultrafine grain structure in the surface layer portion to improve the propagation stopping property of brittle cracks. Since the ultrafine grain structure is thermally unstable, it is essential to devise a manufacturing method so that the ferrite grain size of the ultrafine grain structure does not become coarse or the ultrafine grain structure disappears.

In the present invention, from the results of detailed experiments, from the viewpoint of the relationship with other material characteristics, the ease of production, and the load on the production, it is necessary to have an ultrafine grain structure in the surface layer and a low yield ratio. It was concluded that the following two methods are the most suitable as a production method that is compatible with the introduction of a large proportion of martensite phase.

In the first method, in order to form an ultrafine grain layer on the surface layer, the steel slab is heated to a temperature not lower than the Ac ₃ transformation point and not higher than 1250 ° C., and accumulated in the austenite region at 950 ° C. or lower. After performing rough rolling with a rolling reduction of 10 to 50%, at least two surface layer regions of the outer surface corresponding to 10 to 33% of the thickness of the billet at that stage are treated at a temperature of Ar ₃ transformation point or higher to 2 In the process of starting the cooling at a cooling rate of -40 ° C / s, stopping the cooling at the Ar ₃ transformation point or lower and allowing the heat to be reheated at least once, until the reheat after the last cooling is completed. Finish rolling with a cumulative rolling reduction of 20 to 90% is completed in the meantime, and the surface layer region of the steel material after completion of the rolling is restored to the range of (Ac ₁ transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.). Heat.

The steel material after the end of the recuperation was cooled at a cooling rate of 0.2 to 2 ° C./s, and was cooled to a range of (transformation start temperature (Ar ₃ ) -50 ° C. at the cooling rate) to 500 ° C. rear,
The required martensite phase is formed by cooling to 20-300 ° C at a cooling rate of 5-40 ° C / s.

That is, in order to reduce the yield ratio, the martensite phase is introduced into the matrix phase consisting of the ferrite phase, the bainite phase, and the mixed phase of these structures. For that purpose, an appropriate ferrite / austenite two-phase region is used. After cooling to a certain temperature range, it is rapidly cooled to transform the austenite phase into martensite. With such a manufacturing method, it becomes possible to introduce the required martensite structure without impairing the morphology of the ultrafine grain structure of the surface layer portion.

In the second method, a steel slab is made to have an Ac ₃ transformation point or higher,
After heating to a temperature of 1250 ° C. or lower and performing a rough rolling with a cumulative rolling reduction in the austenite region of 950 ° C. or lower of 10 to 50%, at least corresponding to 10 to 33% of the thickness of the billet at that stage. To start the cooling of the surface layer regions of the two outer surfaces at a cooling rate of 2 to 40 ° C./s from the temperature of the Ar ₃ transformation point or higher, and stop the cooling at the Ar ₃ transformation point or lower to recover the heat 1. In the process of passing more than once, finish rolling with a cumulative rolling reduction of 20 to 90% is completed until the recuperation after the last cooling is completed, and the surface layer region of the steel material after the rolling is completed (Ac ₁ Transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.)
Or re-heat to the range of
Alternatively, by cooling the steel material after the recuperation to 20 to 650 ° C at a cooling rate of 5 to 40 ° C / s, the steel material having an ultrafine grain layer formed on the surface layer is subjected to the following special two-phase heat treatment. Is the way.

That is, when the two-phase region heat treatment for forming martensite is performed by the ordinary heat treatment, the superfine grain structure of the surface layer portion is completely or partially impaired, and therefore it cannot be adopted. By increasing the rate of temperature rise until heating to the phase region temperature and limiting the holding time at the heating temperature to a short time, it is possible to reduce the yield ratio in the structure without impairing the function of the superfine grain surface structure. It becomes possible to introduce an effective martensite phase.

In that case, at a heating rate of 0.1 to 50 ° C./s, (Ac ₁ transformation point + 10 ° C.) to (Ac ₃ transformation point −30 ° C.)
After heating to the range of 1 to 6, the latent time in the temperature range is 1 to 6
It must be 0s. Cooling after heating and holding is preferably quenching for martensite formation, but 0, 5 to 50 ° C / s
It should be in the range of. The specific reasons for limiting the above manufacturing conditions for introducing the martensite phase will be described later.

The above are the requirements for the high-strength steel material for construction which is excellent in fracture resistance of the present invention, but it is necessary to limit the individual chemical components for the reasons described below.

That is, C is contained as an effective component for improving the strength of steel, and if it is less than 0.01%, it is difficult to secure the strength required for structural steel, but if it exceeds 0.15%, it is excessive. Incorporation of the element causes deterioration of the ductile fracture characteristics, resulting in a reduction in the fracture resistance which is the object of the present invention. Further, since toughness and weld cracking resistance are also reduced, 0.01 to 0.15
The range is%.

Next, Si is an element effective as a deoxidizing element and for securing the strength of the base material. However, if the content of Si is less than 0.01%, deoxidation becomes insufficient and it is disadvantageous for securing the strength.
On the contrary, if the content exceeds 1.0%, a coarse oxide is formed and ductility and toughness are deteriorated. Therefore, the range of Si is set to 0.01 to 1.0%.

Further, Mn is an element necessary for securing the strength and toughness of the base material, and is at least 0. It is necessary to contain 1% or more, but the upper limit was set to 2.0% in the range where the material such as toughness and crackability of the welded portion is acceptable.

Al is an element effective in fixing N, which is one of the requirements of the present invention, and an element effective in deoxidizing, refining the γ grain size, etc., but exerts an effect. Is 0.00
It is necessary to contain 3% or more. On the other hand, if it is contained in excess of 0.1%, a coarse oxide is formed and ductility is extremely deteriorated. Therefore, it is necessary to limit the content to 0.003 to 0.1%.

If solid solution N is present, the ductile fracture characteristics deteriorate and the toughness deteriorates after plastic deformation as described above. Therefore, according to the above equation (1) or (2), Al,
It is necessary to contain Ti, Zr, Nb, Ta, V, and B in appropriate amounts. However, it is necessary to limit the total content for the following reasons.

That is, N works effectively for the refinement of γ grains by combining with Al and Ti, so that a small amount works effectively for mechanical properties. Further, it is impossible to industrially completely remove N in steel, and it is not preferable to reduce N more than necessary because it puts an excessive load on the manufacturing process.

Therefore, the lower limit of N is set to 0.00 as an industrially controllable range in which the load on the manufacturing process is allowable.
1% If it is contained excessively, even if the formula (1) or (2) is satisfied, the ductile fracture characteristics and the toughness after plastic deformation may be adversely affected depending on the manufacturing history. Therefore, the upper limit of the allowable range is 0. 0.006%.

As described above, since P deteriorates the ductility of the ferrite matrix, it is necessary to limit the content of P in order to improve the plastic deformability, the generation of ductile cracks, and the growth characteristics. The smaller the amount of P is, the more preferable it is. However, reducing the amount of P puts a load on the refining process and causes a decrease in productivity and an increase in cost.
The lower limit of is set to 0.01% or less based on the experimental result.
That is, although the ductility characteristic deteriorates as the P content increases,
When it exceeds 0.01, the degree becomes remarkable. P amount is 0.
If it is less than 01%, the adverse effect of P is small. Therefore, in the present invention, the amount of P as an impurity is set to 0.01%.
Limited to:

However, when it is used in a structure where local plastic deformation in the segregation portion or deterioration of ductile fracture characteristics has an effect, the P content is 0.
It is more preferable to limit it to 007% or less.

As for S, as described above, since MnS is formed, the ductile fracture characteristic is deteriorated. In particular, it deteriorates the propagation characteristics of ductile cracks. Since the characteristic of ductile fracture is deteriorated under the condition of a large amount of solid solution P and N, deterioration of the ductile crack propagation characteristic due to S has a great influence on the plastic deformability and the ductile fracture characteristic of the entire steel material. , S must be extremely reduced to about 0.001% or less.

However, if the amount of P and N is reduced and the solid solution N is fixed by the nitride forming element as in the present invention,
Since the allowable amount of S widens because the resistance until the occurrence of ductile fracture becomes large, the present invention limits S as an impurity to 0.01% or less based on the experimental results.

As for O, as described above,
Is preferably reduced as much as possible in order to form inclusions harmful to ductility, but like S, if the solid solution P and N are reduced, ductility of the base material is secured to some extent, so solid solution P, N
The allowable amount is higher than the case where the reduction of
Based on the experimental results, it is necessary to limit it to 0.006% or less.

Ti, Zr, Nb, Ta, V, and B are mainly used for N-fixing, and optionally one or more of them are selectively contained, but the individual elements are also selected for the reasons described below. It is necessary to limit the amount of the components.

Ti is an element effective for fixing N, and contributes to the improvement of the base metal strength by precipitation strengthening.
Although it is an element effective for making γ grains fine by forming TiN, it is necessary to contain 0.003% or more in order to exert the effect. On the other hand, if it exceeds 0.02%, coarse precipitates and inclusions are formed to deteriorate toughness and ductility, so the upper limit is made 0.02%.

Zr is also an element that forms a nitride and is effective in fixing N. However, in order to exert its effect, Zr.
It is necessary to contain 003% or more. On the other hand, if it exceeds 0.10%, as with Ti, coarse precipitates and inclusions are formed to deteriorate the toughness and ductility, so 0.003 to 0.10.
Limit to the range of%.

Nb is also an element effective for fixing N, but if contained in excess, the toughness deteriorates due to precipitation embrittlement. Therefore,
As a range that can exert the effect without causing deterioration of toughness,
It is limited to the range of 0.002 to 0.05%.

Ta is also an element effective in fixing N, but in order to exert the effect, it is necessary to contain 0.005% or more. On the other hand, if it exceeds 0.20%, precipitation embrittlement, coarse precipitates, and toughness deterioration due to inclusions occur, so the upper limit is made 0.20%.

V is also an element effective for fixing N by forming VN, but as with Nb, if contained in excess, the toughness deteriorates due to precipitation embrittlement. Therefore, the range in which the effect can be exhibited without causing deterioration of toughness is limited to the range of 0.005 to 0.20%.

Since B is a trace amount and is surely bound to N, it is an element effective for N fixing, and 0.0 is necessary for exerting the effect.
002% or more is required. On the other hand, if it is contained in excess of 0.003%, the BN becomes coarse, and the ductility and toughness are adversely affected. Moreover, since the weldability is also deteriorated, the upper limit is set to 0.
003%.

In addition to the above, in order to increase the strength of the base material and to secure the toughness in accordance with the desired strength level, C is added as necessary.
One or more of r, Ni, Mo, Cu and W may be contained. First, Cr and Mo are both effective elements for 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 2.0%, Since the toughness and the weldability tend to deteriorate, the respective ranges are 0.01 to 2.0%.

Further, 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 large, the strength and toughness are improved, but the effect is saturated even if added in excess of 4.0%, but the weldability deteriorates, so the upper limit is made 4.0%.

Next, Cu has almost the same effect as Ni, but if it exceeds 2.0%, a problem occurs in hot workability.
It is limited to the range of 0.01 to 2.0%. W is effective in increasing the strength of the base metal due to solid solution strengthening and precipitation strengthening, but 0.01% or more is necessary to exert the effect. on the other hand,
If the content is excessively over 2.0%, the toughness will be significantly deteriorated, so the upper limit is made 2.0%.

Further, in order to improve ductility and joint toughness, one or more of Mg, Ca and REM may be contained, if necessary. Mg, Ca, REM
All of these are effective in improving the ductility characteristics by suppressing the expansion of the sulfide during hot rolling. It also works effectively to improve the joint toughness by refining the oxide. The lower limit content for exhibiting the effect is 0.0005% for Mg and Ca, and 0.005% for REM. On the other hand, if it is contained excessively, coarsening of sulfides and oxides occurs, leading to deterioration of ductility and toughness. Therefore, the upper limits are set to 0.01% for Mg and Ca, and 0.10% for REM.

Next, the reasons for limiting the production of the low yield ratio high temperature steel for construction having excellent fracture resistance of the present invention will be described. In the steel having a limited chemical composition for the above reasons, in order to improve the brittle crack propagation arresting property, in the surface layer portions of at least two surfaces of the steel material, an ultrafine grain structure having an average ferrite grain size of 3 μm or less is formed from the surface layer. Plate thickness 10-33
% Must be present over the thickness. The superfine grain surface layer having the characteristics limited in the present invention can be formed by limiting the production conditions as shown below.

During hot rolling of a steel slab, a region of an appropriate thickness of the surface layer portion is once cooled to a temperature lower than the Ar ₃ transformation point during or during hot rolling by means such as water cooling. After making a temperature difference with the inside, when hot rolling is further performed from the state where the temperature difference remains, the area once cooled to a temperature lower than the Ar ₃ transformation point is reheated and processed in that process to produce ferrite. Become the main organization.

Therefore, the surface layer portion having the ferrite main structure is subjected to processing while being reheated by the sensible heat inside, and by optimizing the processing conditions during this reheating, the ferrite crystal grains in the surface layer portion are Remarkably fine-grained.
Therefore, the ratio of the surface ultrafine grain layer in the final steel product is
When the surface layer is once cooled, it almost coincides with the ratio of the region reduced to the Ar ₃ transformation point.

In the above hot rolling step, ultrafine graining can be achieved by satisfying the following conditions. First, the steel slab is reheated to the austenite region, and the temperature in this case is preferably the Ac ₃ transformation point or higher and 1250 ° C. or lower. That is, below the Ac ₃ transformation point, the austenite single phase is not formed and the ferrite phase remains. If the ferrite phase remains, a uniform ultrafine grain structure cannot be formed in the surface layer regardless of the subsequent steps. .

Further, since the inside is also processed in the two-phase region, there is a problem that the anisotropy of the steel material increases. On the other hand, if the temperature exceeds 1250 ° C, the heated austenite grain size becomes extremely coarse,
The grain size cannot be reduced even by the subsequent rolling, and the toughness at the center of the plate thickness cannot be secured. Therefore, in the present invention, the heating temperature of the billet is limited to the Ac ₃ transformation point to 1250 ° C.

After heating the steel slab, rolling with a cumulative reduction of 10 to 50% is performed in the austenite region of 950 ° C. or lower. This is a refinement of the austenite grain size before transformation,
This is because the surface layer has an ultrafine grain structure in a later step, and the toughness of the internal normal structure is ensured by setting the crystal grain size at the center of the plate thickness to 30 μm or less. The term "substantial refinement of austenite grains" means refinement of recrystallized austenite as well as expansion of austenite grains by non-recrystallization rolling.

The reduction in the austenite region at a low temperature is effective for substantially refining austenite, but the reduction above 950 ° C. is not effective for refining austenite. Therefore, in the present invention, reduction at a temperature below 950 ° C. Limit the rolling reduction. If the rolling reduction at 950 ° C. or less is less than 10%, the effect of processing is insufficient, and therefore there is no effect on the refinement of austenite.

The larger the rolling reduction in the austenite region of 950 ° C. or lower, the more effective the refinement of austenite is. However, if the rolling reduction exceeds 50%, the effect tends to be saturated, and the rolling reduction exceeds 50%. When it becomes larger, it is effective for making the austenite finer, but it becomes impossible to secure the reduction rate in the recuperating process which is essential for making the ferrite of the surface layer later to be finer. The upper limit of the rolling reduction below is 50%.

Note that the reduction at a temperature higher than 950 ° C. has a small effect on the refinement of austenite, but does not adversely affect the material as long as the required reduction rate in the subsequent reheating step can be secured. Therefore, when the initial slab thickness is large, the processing may be performed at a temperature higher than 950 ° C. as necessary.

Under the above conditions, the austenite grains were sufficiently refined and unrecrystallized region rolling was performed, and then the surface layer portion of the steel material to be an ultrafine grain layer was cooled by means of water cooling or the like, and the steel material was water-cooled. The area of each surface layer portion corresponding to 10 to 33% of the sheet thickness at the time of the previous hot rolling is cooled to the Ar ₃ transformation point or lower, and a temperature difference is provided between the surface layer portion and the inside.

At that time, the cooling rate of the region of each surface layer portion corresponding to 10 to 33% of the plate thickness at the time of hot rolling before water cooling of the steel material needs to be 2 ° C./s or more. This is because if the cooling rate is less than 2 ° C / s, the transformation structure after cooling becomes coarse even if the austenite is refined by hot rolling before cooling.
This is because it is difficult to obtain a uniform ultrafine ferrite structure by subsequent rolling during recuperation.

A higher cooling rate is preferable from the viewpoint of micronization of the structure, but the effect is saturated even if the material is rapidly cooled above 40 ° C./s, and unnecessary cooling is necessary to maintain the shape of the steel sheet. Therefore, the upper limit is set to 40 ° C./s.

Further, the cooling after the rolling at 950 ° C. or lower is started from the Ar ₃ transformation point or higher. this is,
This is because the superfine grain surface layer is uniformly formed by cooling the single-phase austenite. That is, when the surface layer portion is below the Ar ₃ transformation point before forced cooling, ferrite is partly coarsely formed and superfine grain formation is hindered at that portion.

For the above reasons, the region of each surface layer portion corresponding to 10 to 33% of the plate thickness at the time of hot rolling of the steel material before cooling is cooled at an Ar ₃ transformation point at a cooling rate of 2 to 40 ° C./s. When cooling to the following and then performing finish rolling, the sensible heat of the inside and / or the heating from the outside may be used to reduce the plate thickness to 10
By rolling the region of each surface layer portion corresponding to ˜33% while the temperature is rising, the structure of the region becomes ultra-fine and can contribute to the improvement of the brittle crack propagation arresting property.

As will be described later, the processing in the recuperation process is 1
The recuperation temperature after rolling in the recuperation process after the last cooling may be (Ac ₁ transformation point-
50 ° C.) to (Ac ₃ transformation point + 50 ° C.). That is, the final recuperation temperature is (Ac ₁ transformation point −50
If the temperature is lower than (° C.), recovery and recrystallization of the worked ferrite after working are not sufficient, so that ultrafine graining is insufficient and brittle crack propagation arresting property is not improved.

On the other hand, the final recuperation temperature is (Ac ₁ transformation point +
If the temperature is higher than 50 ° C, part of the ultrafine-grained ferrite disappears due to the reverse transformation to austenite again by working, and the ratio becomes so large that it cannot be ignored, and the toughness and brittle crack propagation arresting characteristics are increased. Spoil. Therefore, in the present invention, the recuperation temperature after rolling in the recuperation process after the last cooling is limited to the range of (Ac ₁ transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.).

The above-described processing steps during cooling to below the Ar ₃ transformation point and during the reheating may be carried out once, but the effect is superposed by repeating it a plurality of times, so the desired microstructure can be obtained even if it is repeated twice or more. Is possible.

In this case, even if the maximum temperature or the minimum temperature of each recuperation step is arbitrary, ultrafine particles are obtained according to the temperature conditions of the present invention. However, it is preferable that the upper limit of the recuperation temperature on the way is (Ac ₃ transformation point + 100 ° C.) or less, because the grain refining effect is surely superposed.

If the cumulative reduction ratio of finish rolling as the rolling from the initial cooling to the final recuperation is large, an ultrafine grain structure can be obtained uniformly and stably. For that purpose, the cumulative rolling reduction of finish rolling must be at least 20%. The larger the rolling reduction is, the more advantageous the ultrafine graining is, but the rolling reduction is 9
Rolling exceeding 0% is not preferable because the effect is saturated and productivity is extremely hindered. Therefore, in the present invention, the cumulative reduction rate of finish rolling is limited to 20 to 90%.

By the manufacturing method according to the above-mentioned limiting conditions,
It is possible to add an ultrafine grained layer to the surface layer portion, but in order to further reduce the yield ratio, it is necessary to limit the cooling conditions after rolling or the heat treatment conditions after steel material production as shown below. .

In the method of introducing the martensite phase in the cooling stage after the end of the final recuperation, the steel material after the end of recuperation is cooled at a cooling rate of 0.2 to 2 ° C./s (transformation at the cooling rate). after cooling to a range of start temperature _{(Ar 3) -50 ℃) ~500} ℃, 20~300 at a cooling rate of 5 to 40 ° C. / s
Cool to ° C.

First, the steel material after the completion of recuperation is cooled to the temperature of the two-phase region. If the cooling rate at that time is less than 0.2 ° C./s, the cooling rate is too slow, so ferrite or bainite formed by transformation, Or, because the matrix structure, which is a mixed phase of these, becomes coarse, causing deterioration of toughness,
It is not preferable because the crystal grain size of the ultrafine grain layer formed in the previous stage may become coarse and the brittle crack propagation arresting property may be deteriorated.

If the cooling rate is higher than 2 ° C./s, the transformation start temperature becomes too low, which makes it difficult to form a two-phase structure of a transformation phase and an austenite phase, or a matrix and martensite. It is not preferable because the difference in hardness between the two becomes small and the yield ratio cannot be lowered.

For the above reasons, in the present invention, from the last recuperation (the transformation start temperature (A
r ₃₎ -50 ℃) to limit the scope of the cooling rate to to 500 ° C. in 0.2 to 2 ° C. / s.

After the completion of the recuperation, the temperature of the two-phase region was cooled at 0.2 to 2 ° C./s to optimize the ratio of the mother phase generated by the transformation and the untransformed austenite phase. Quench to transform to martensite phase. At that time, the temperature at which the cooling at 0.2 to 2 ° C./s is stopped is (transformation start temperature (Ar ₃ ) −5 at the cooling rate).
It should be in the range of 0 ° C) to 500 ° C.

In order to stably form the martensite phase of 10 to 60% necessary for lowering the yield ratio in the structure, it is necessary to concentrate C in austenite to a certain level or more.
For that purpose, it is necessary to set the temperature to 50 ° C. or lower than the transformation start temperature (Ar ₃ transformation point) at the cooling rate until entering the two-phase region. However, if this temperature becomes too low, transformation occurs before the subsequent quenching stage, and a bainite phase with a small hardness difference from the parent phase may be generated instead of a hard martensite with a high concentration of C. Will be more likely.

According to the experimental results, the lower limit temperature for ensuring the proportion of martensite of 10 to 60% is 500 ° C. Therefore, the cooling stop temperature before quenching in the present invention, (the transformation start temperature in cooling rate (Ar ₃₎ -5
(0 ° C) to 500 ° C.

The untransformed austenite is transformed into a martensite phase by rapid cooling from Ar ₃ -50 to 500 ° C. The higher the cooling rate is, the more advantageous for the martensite transformation. Considering the transformation from austenite, the lower limit of the cooling rate is 5
It should be ℃ / s.

Further, although increasing the cooling rate is advantageous for martensitic transformation, it causes a rise in manufacturing cost, and there is a problem that residual stress remains in the steel material to cause deformation of the steel material. The upper limit of the cooling rate is 40 ° C./s, which is sufficient and does not cause the above problems.

After cooling at 5 to 40 ° C./s to cause martensitic transformation, cooling can be stopped midway for the purpose of reducing residual stress and improving the material. As the quenching stop temperature after martensitic transformation, cooling to a low temperature exceeding 20 ° C. is meaningless since it has no effect on the properties of martensite, and when quenching is stopped at a high temperature of more than 300 ° C., The martensite transformation is not yet completed, and untransformed austenite may transform into the bainite phase, and the required amount of martensite may not be secured. Therefore, the quench stop temperature is limited to the range of 20 to 300 ° C.

When the steel material after rolling / cooling is subjected to heat treatment for reheating to produce a martensite phase, the steel material after the end of the recuperation is allowed to cool, or the steel material after the end of the reheat is in a range of 5-40.
After cooling to 20 to 650 ° C. at a cooling rate of C / s, the temperature is further increased to 0.1 to 50 ° C./s (Ac ₁ transformation point + 10
℃) ~ (Ac ₃ transformation point -30 ℃) after heating in the range,
After holding for 1 to 60 s in the temperature range, 0.5 to 50 ° C / s
Cool with.

After the hot rolling including the rolling in the recuperating step according to the conditions of the present invention is performed to form the superfine grain structure in the surface layer portion, the cooling rate is such that the subsequent grain growth can be suppressed. It may be cooled to a temperature at which the transformation is substantially completed.

When the plate thickness of the thickest cross section of the steel material is 100 mm or less, even cooling is sufficient. In addition, it is also possible to quench rapidly for the purpose of strength adjustment and shortening of manufacturing time,
In the present invention, the cooling condition in that case is to cool to 20 to 650 ° C at a cooling rate of 5 to 40 ° C / s.

Limiting the cooling rate to 5 to 40 ° C./s is as follows.
This is because if it is less than 5 ° C / s, there is no effect on strength adjustment.
This is because if the temperature exceeds ℃ / s, the strength increasing effect and the structure controlling effect are saturated, but the deformation and residual stress of the steel material tend to be large, and it is meaningless to further increase the cooling rate in practice.

The temperature at which the rapid cooling is stopped at 5 to 40 ° C./s is in the range of 20 to 650 ° C. This has no effect on the material and structure control even if the rapid cooling is performed to less than 20 ° C. On the other hand, there is a possibility that the manufacturing cost may rise and the shape of the steel material may deteriorate.
If it is over, the transformation near the plate thickness center is still in progress, so the desired material cannot be obtained due to the coarsening of the structure and the increase of high temperature transformation products, and the intention of rapid cooling for material control is completely lost. This is because it will end up.

The steel material produced by the above method is reheated to the two-phase region to generate the required amount of martensite phase. The requirements for the heat treatment are (A
After heating in the range of (c ₁ transformation point + 10 ° C.) to (Ac ₃ transformation point −30 ° C.) and then maintaining the temperature range for 1 to 60 s,
It is to cool at 0.5 to 50 ° C./s.

A problem in performing the two-phase region heat treatment is how to preserve the ultrafine grain structure of the surface layer portion formed in the rolling step. The superfine grain structure of the surface layer is a structure formed by a specially devised thermal history, so that not only the temperature exceeding the transformation temperature but also the tempering treatment at a high temperature causes recrystallization, grain growth, etc. This increases the possibility of damaging the peculiar ultrafine grain structure.

As the heat treatment for introducing martensite while maintaining the ultrafine grain structure, a two-phase region heat treatment of rapid heating and holding for a short time is essential. That is, by rapidly heating, it is possible for the ultrafine grain structure to reach the two-phase region temperature before it changes with its heat as a driving force. Similarly, by holding for a short time, the ultrafine grain structure in the holding stage It becomes possible to suppress the grain growth of the tissue.

In this case, the temperature rising rate needs to be in the range of 0.1 to 50 ° C./s. If the heating rate is less than 0.1 ° C./s, there is no effect of rapid heating, and it is difficult to suppress grain growth in the ultrafine grain portion. On the other hand, if it exceeds 50 ° C./s, although it is effective in suppressing the grain growth of the ultrafine grain portion, the holding temperature easily overshoots and industrially stable control becomes difficult. Therefore, the upper limit of the present invention is 50. Limited to ° C / s.

It is preferable to control the heating rate from the beginning to the holding temperature within this temperature increase rate range.
If the average temperature rising rate from the temperature of ℃ to the holding temperature is within the range of the present invention, the two-phase region heat treatment can be performed without damaging the superfine grain structure of the surface layer portion.

By optimizing the heating temperature within the limited range of the above heating rate and the holding time described later, the ratio of the martensite phase in the steel material after the heat treatment is suitable for lowering the yield ratio.
It is controlled to be in the range of 0 to 60%. For that purpose, (Ac ₁ transformation point + 10 ° C.) to (Ac ₃ transformation point −30)
It is necessary to heat to a two-phase region temperature in the range of (° C).

If the heating temperature is lower than (Ac ₁ transformation point + 10 ° C.), the proportion of the austenite phase formed during heating is small, so that the proportion of the martensite phase formed by transformation during cooling cannot be 10% or more. .

On the contrary, the heating temperature is (Ac ₃ transformation point −30 ° C.)
If it is over, the concentration of C in the austenite phase formed at the time of heating is not sufficient and the hardenability of austenite, which does not depend on the chemical composition, cannot be ensured. This is because the transformation is not assured and it becomes difficult to stably obtain the required amount of martensite, and there is an increased risk that the morphology of the superfine grain structure in the surface layer portion will collapse when the heating temperature rises.

Therefore, in the present invention, the heating rate is 0.1 to 50 ° C./s and the holding time at the heating temperature is 1 to 60.
If it is assumed that s is s, the heating temperature for the two-phase heat treatment is (Ac) so that the required amount of martensite can be stably secured and the morphology of the superfine grain structure of the surface layer portion is not impaired. ₁ transformation point + 10 ° C) ~ (Ac ₃ transformation point -30 ° C)
It is limited to the range of.

The reason why the holding time at the heating temperature is limited to 1 to 60 s is that the morphology of the superfine grain structure in the surface layer portion is not impaired as in the case of increasing the temperature rising rate. If the holding time is less than 1 s, industrial control is difficult, and if it exceeds 60 s, recrystallization and grain growth of the superfine grain structure of the surface layer portion start.

It should be noted that increasing the rate of temperature rise and limiting the holding time at the heating temperature to a short time are effective in preserving the structure of the surface layer portion, and at the same time, the martensite phase during the two-phase region heat treatment is performed. It also has a complementary effect on the refinement of, and is effective for improving toughness.

The cooling condition after holding for 1 to 60 s at (Ac ₁ transformation point + 10 ° C.) to (Ac ₃ transformation point −30 ° C.) should be within a range where a necessary amount of martensite phase is formed during cooling transformation. Good. In the present invention, the cooling rate is 0.5.
If it is less than ° C / s, the martensite phase is not surely formed, and the faster the cooling rate is, the more advantageous it is.
If the temperature exceeds ℃ / s, the effect will be saturated for the formation of martensite phase during the two-phase heat treatment, but there will also be a demerit in terms of the shape and cost of the steel material.
It is limited to the range of.

The above-mentioned manufacturing method in which the ultrafine grained layer according to claim 5 is formed and then immediately rapidly cooled from the two-phase region temperature, or the rapid heating and the short-term holding described in claim 6 are characterized. It is also possible to further temper the steel material manufactured by the manufacturing method by the two-phase heat treatment, for the purpose of adjusting strength, improving toughness, and improving shape. In that case, it is an essential condition that the ultrafine grain structure formed in the surface layer portion is not damaged.

In the present invention, the tempering temperature is 450 to 650 ° C.
This is because the effect of tempering is not clear below 450 ° C., and the morphology of the superfine grain structure in the surface layer portion may be impaired above 650 ° C. It should be noted that the heating holding time for tempering is arbitrary within the tempering temperature range, but from the viewpoint of preserving the superfine grain structure of the surface layer portion, the holding time is preferably within 5 hours.

[0149]

[Examples] Using test steels having the chemical compositions shown in Table 1, Table 2
For a thick steel plate having a plate thickness of 50 mm manufactured under the manufacturing conditions shown in (1), as-manufactured strength and toughness (fracture transition temperature vTrs) of a base metal after a strain of 10% by a Charpy test,
Brittle crack propagation arrest property (Kca value of 4 by ESSO test)
00 kgf ・ mm ^-3/2 ), the limit of ductile fracture occurrence C
Table 3 shows the TOD value (δi), welded joint properties (Charpy properties as-welded (average value of absorbed energy at −20 ° C.), as-welded and δi after 10% strain).

[0150]

[Table 1]

[0151]

[Table 2]

[0152]

[Table 3]

[0153]

[Table 4]

[0154]

[Table 5]

[0155]

[Table 6]

[0156]

[Table 7]

[0157]

[Table 8]

[0158]

[Table 9]

[0159]

[Table 10]

[0160]

[Table 11]

[0161]

[Table 12]

[0162]

[Table 13]

[0163]

[Table 14]

[0164]

[Table 15]

[0165]

[Table 16]

The tensile properties of the base material were measured by a round bar test piece having a parallel part diameter of 6 mm and an inter-rating distance of 25 mm, which was taken from the t / 4 part of the plate thickness so that the test direction was perpendicular to the rolling direction. Carried out. The Charpy impact characteristics of the base material were also sampled at the same position and direction as the tensile test piece, and the fracture surface transition temperature (vTrs) was determined.

Critical CTOD value (δi) for occurrence of ductile fracture
The test was performed using a three-point bending test piece with a fatigue notch that was taken from the center of the plate thickness so that the longitudinal direction of the test piece was perpendicular to the rolling direction.

The welding conditions were submerged arc welding of one layer on both sides. Welding was performed by adjusting the welding heat input so as to fall within the range of about 190 to 200 kJ / cm. The 2mmV notch Charpy impact test piece of the joint and the test piece for measuring δi are
Set the position 7 mm below the surface to be the center of the test piece,
The sample was sampled so that the position 1 mm into the HAZ side from the boundary between the weld metal and the HAZ (fusion part: FL) would be the notch position. The tensile properties, the base metal, and the δi of the joint were all measured at room temperature. Regarding the pre-strain test, a plate-shaped test piece was cut out from a steel sheet and applied with a tensile strain of 10%, and then each test piece was sampled and subjected to a characteristic investigation.

As shown in Tables 1 and 2, steel numbers A1 to A1
Steel sheet No. 6 has a chemical composition within the scope of the present invention and a surface superfine grain structure, and exhibits a low yield ratio of less than 80%,
Temperature Kca value determined by ESSO test is indicative of the propagation stop characteristics of brittle cracks becomes 400 kgf / mm ^-3/2 is not only very good, even after applying a large strain of 10% In addition to the elongation required for a normal round bar tensile test, which exhibits very little deterioration, δi, which indicates the characteristic of ductile fracture in the presence of cracks, maintains a good value regardless of whether or not strain is applied. The Charpy characteristics of the welded joint also have the characteristics necessary for safe use in structures such as buildings and bridges, and the δi of the joint, like the base metal, is sufficiently high even after straining. However, the steel material produced by the present invention, even when used in a structure that undergoes large and repeated plastic strain due to a large earthquake during use,
It is clear that it is a low yield ratio steel material with a high level of safety never before seen.

On the other hand, steel numbers B1 to B10 are comparative examples,
One of the properties shown in Table 3 is inferior to the steel of the present invention because it does not satisfy the requirements of the present invention. That is,
Steel No. B1 has an excessive total N content, so ESS before straining
The O property is also inferior to that of the steel of the present invention, but especially E after straining
SSO characteristics and δi are inferior.

Steel No. B2 has an excessive amount of total N and an insufficient amount of Al, Ti, and Nb for fixing the solid solution N. Therefore, the brittle crack propagation arresting property after strain application and the value of δi Is deteriorated as compared with the steel of the present invention. Steel No. B3 is within the chemical composition range of the present invention as the total amount of N, but because the fixation of N is insufficient, that is, the value of the formula (1) is a positive value, the deterioration of the material due to the strain application is caused. large.

Steel No. B4 is Al most effective for fixing solid solution N.
Since the content of N is insufficient, N is not fixed enough,
Degradation of ESSO characteristics and δi due to application of strain is remarkable. Steel No. B5 has an excessive amount of P, so that the ductile fracture characteristics and ESSO characteristics are relatively low even before the strain is applied, and further, the ductile fracture characteristics and the ESSO characteristics after the strain are applied are significantly reduced.

Since steel No. B6 has an excessive amount of S, the ductility characteristics (elongation, δi) are significantly inferior to those of the steel of the present invention both before and after strain is imparted. Since steel No. B7 has an excessive amount of C, the ductile fracture characteristics and ESSO characteristics before and after the application of strain are relatively low, and the toughness of the welded joint is remarkably poor.

Steel No. B8 is within the scope of the present invention as a chemical component, but since it does not have a superfine grain structure in the surface layer portion,
The ESSO characteristics are significantly deteriorated before and after the strain is applied. Steel No. 9 has a fine-grained structure in the surface layer as compared with the center, but its grain size does not satisfy the requirements of the present invention and is coarse. Cannot be obtained before and after applying strain.

In Steel No. 10, since the thickness of the superfine grain surface layer is insufficient, sufficient brittle crack propagation arresting properties cannot be obtained before and after applying strain. Steel No. B11 does not undergo accelerated cooling from the temperature of the two-phase region after rolling and is not subjected to rapid heating and tempering treatment to the two-phase region, so the martensite ratio in the structure is too small and the yield ratio is for construction. It is insufficient as a low yield ratio steel.

On the contrary, with steel No. B12, the heating temperature for two-phase tempering is too high and martensite is surplus, so that both the Charpy property and the brittle crack propagation stopping property are significantly deteriorated.
Steel No. B13 does not satisfy the requirements for the two-phase region tempering conditions of rapid heating to the two-phase region and holding for a short time, which are features of the present invention. That is, since the temperature rising rate of the two-phase region tempering is slow and the holding time is excessive, the morphology of the superfine grain layer of the surface layer portion once formed collapses, and since the average grain size becomes coarse, even before straining, No improvement in brittle crack propagation arresting property.

Steel No. B14 is particularly inferior in ductile fracture characteristics because the O content is excessive. Steel No. B15 has not been rolled in the γ region before cooling the surface layer, so the average grain size at the center of the plate thickness becomes coarse, the internal toughness is poor, and brittle crack propagation stops due to plastic deformation. The characteristic deterioration is remarkable.

From the above examples, according to the present invention, not only before applying the pre-strain, but also after applying the pre-strain of 10% which is assumed to be subjected to a large deformation such as a large earthquake,
It is clear that it is possible to obtain a steel material with very good Charpy properties, brittle crack propagation arrest properties and ductile fracture properties (drawing value, δi).

[0179]

INDUSTRIAL APPLICABILITY The present invention has very little deterioration of the material due to plastic strain even when it is subjected to large and repeated plastic strain due to a large earthquake or the like during use, and brittle cracks are generated even after plastic deformation. It suppresses the initiation and propagation of ductile cracks that make it easier, and even if a fracture occurs, it can stop the brittle cracks. This is made possible in a normal steel material manufacturing process without using, and its industrial effect is extremely large.

─────────────────────────────────────────────────── ─── Continuation of front page (51) Int.Cl. ⁷ Identification code FI C22C 38/58 C22C 38/58 (56) Reference JP-A 61-235534 (JP, A) JP-A 64-47815 (JP , A) JP-A-7-126798 (JP, A) (58) Fields investigated (Int.Cl. ⁷ , DB name) C22C 38/00-38/60 C21D 8/00

Claims (7)

(57) [Claims] 1. In mass %, C: 0.01 to 0.15% Si: 0.01 to 1.0% Mn: 0.1 to 2.0% Al: 0.003 to 0.1% N : 0.001 to 0.006%, and N (%)-Al (%) / 3.0 ≦ 0, and the content of P, S, O as impurities is P: 0.01% Hereinafter, S: 0.01% or less, O: 0.006% or less, a steel material composed of the balance iron and unavoidable impurities, the average crystal grain size of the center part of the plate thickness is 30 μm or less, and further occupies the steel material volume. The martensite ratio is 10 to 60%, and further 10 to 33% of the total thickness from the surface layer with respect to at least two outer surfaces of the outer surfaces constituting the steel material.
A low yield ratio high tensile steel material for construction having an excellent fracture resistance, characterized by having an ultrafine grain structure with an average ferrite grain size within the range of 3 μm or less. 2. In mass %, Ti: 0.003 to 0.020% Zr: 0.003 to 0.10% Nb: 0.002 to 0.050% Ta: 0.005 to 0.20% V : 0.005 to 0.20% B: 0.0002 to 0.003% of 1 type or 2 types or more, N (%)-Al (%) / 3.0-Ti (%) / 3.4-Zr (%) / 6.5-Nb (%) / 13.2-Ta
(%) / 25.8-V (%) / 10.9-B (%) / 2.0 ≦ 0. The low yield ratio, high tensile steel material for construction with excellent fracture resistance according to claim 1. 3. In mass %, Cr: 0.01 to 2.0% Mo: 0.01 to 2.0% Ni: 0.01 to 4.0% Cu: 0.01 to 2.0% W : 0.01 to 2.0% of 1 type or 2 types or more are contained, The low yield ratio high tensile steel material for construction excellent in fracture resistance of Claim 1 or 2 characterized by the above-mentioned. 4. In mass %, Mg: 0.0005-0.01% Ca: 0.0005-0.01% REM: 0.005-0.10% One or more kinds are contained. The low-yield-ratio, high-strength steel material for construction having excellent fracture resistance according to any one of claims 1 to 3. 5. A steel slab having the composition according to any one of claims 1 to 4 is heated to a temperature not lower than the Ac ₃ transformation point and not higher than 1250 ° C., and the cumulative rolling reduction in the austenite region of not higher than 950 ° C. is 10. After carrying out rough rolling of ˜50%, at least two surface layer regions of the outer surface corresponding to 10 to 33% of the thickness of the billet at that stage are heated at a temperature of Ar ₃ transformation point or higher to 2 to 40 ° C.
In the process of starting the cooling at a cooling rate of / s and stopping the cooling at the Ar ₃ transformation point or lower to reheat the heat at least once, the reheating after the last cooling from the start of the cooling is completed. Finish rolling with a cumulative rolling reduction of 20 to 90% is completed in the range of (Ac ₁ transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.) in the surface layer region of the steel material after the rolling. The steel material after the recuperation is 0.2-2 ℃ / s
At the cooling rate of (the transformation start temperature (Ar
₃ ) After cooling in the range of -50 ° C) to 500 ° C, 5-4
A low yield ratio high for construction excellent in fracture resistance, characterized in that the steel material according to any one of claims 1 to 4 is manufactured by cooling to 20 to 300 ° C at a cooling rate of 0 ° C / s. Method of manufacturing tensile steel. 6. A steel slab having the composition according to any one of claims 1 to 4 is heated to a temperature of Ac ₃ transformation point or higher and 1250 ° C. or lower, and a cumulative rolling reduction in an austenite region of 950 ° C. or lower is 10. After carrying out rough rolling of ˜50%, at least two surface layer regions of the outer surface corresponding to 10 to 33% of the thickness of the billet at that stage are heated at a temperature of Ar ₃ transformation point or higher to 2 to 40 ° C.
In the process of starting the cooling at a cooling rate of / s and stopping the cooling at the Ar ₃ transformation point or lower to reheat the heat at least once, the reheating after the last cooling from the start of the cooling is completed. Finish rolling with a cumulative rolling reduction of 20 to 90% is completed in the range of (Ac ₁ transformation point −50 ° C.) to (Ac ₃ transformation point + 50 ° C.) in the surface layer region of the steel material after the rolling. After the recuperation, the steel after the recuperation is left to cool, or the steel after the recuperation is completed at 20 to 6 at a cooling rate of 5 to 40 ° C / s.
After cooling to 50 ° C., the temperature is further increased from 0.1 to 50 ° C./s at a rate of (Ac ₁ transformation point + 10 ° C.) to (Ac ₃ transformation point −3).
0 ° C.) range, after holding for 1 to 60 s in the temperature range, it is subjected to a two-phase zone heat treatment of cooling at 0.5 to 50 ° C./s. A method for producing a high yielding steel material having a low yield ratio for construction, which is excellent in fracture resistance and is characterized by producing the above steel material. 7. The method for producing a high yield steel material having a low yield ratio for construction excellent in fracture resistance according to claim 5, wherein tempering is performed at 450 to 650 ° C.