JP2014029019A - Method for producing steel sheet for large heat input welding excellent in brittle crack arrest property - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a method for producing a thick steel plate having a sheet thickness of 40 mm or more, suitable for large heat input welding and excellent in brittle crack arrest properties.SOLUTION: A steel material contains specific amounts of C, Si, Mn, P, S, Al, Nb, Ti, N, B, Ca and the balance of Fe with inevitable impurities and has a Ceq expressed by expression (1) of 0.33 to 0.45, the mass% of Ca, S and O in the steel satisfying a specific expression and mass% of Ti, B and N in the steel satisfying another specific expression. The steel material is heated to 1,000°C or more, rolled at an austenite recrystallization region, subjected to accelerated cooling to a non-recrystallization temperature range, then rolled at cumulative rolling reduction of 40% or more at an austenite non-recrystallization temperature range, and is subjected to accelerated cooling from Artransformation point or more to a temperature range of 500°C or less. Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 (1), each symbol of the elements is the content of the content (mass%) of each element.

Description

  The present invention is a method for producing a steel plate having a thickness of 40 mm or more used for various large steel structures such as ships, marine structures, low-temperature storage tanks, and construction / civil engineering structures, and is particularly suitable for large heat input welding. In addition, the present invention relates to a suitable method for producing a thick steel plate having excellent brittle crack propagation stopping characteristics.

  Steel structures used in the field of various large steel structures such as ships, offshore structures, low-temperature storage tanks, and construction / civil engineering structures have the required characteristics of strength and toughness from the viewpoint of ensuring their safety. Satisfaction is required.

  Moreover, in recent years, in addition to the above-mentioned required characteristics generally attached to steel sheets, there is an increasing demand for steel sheets having brittle fracture propagation stop characteristics for the purpose of preventing an increase in the destruction accidents of steel structures. . As a background to this, for example, in the marine field, the amount of construction of larger vessels tends to increase, so that the thickness of the applied steel sheet is increased and higher strength is aimed at, and the marine structure In the field, the construction of drilling rigs in cold regions is expected.

  Brittle fracture of steel sheets is a type of steel sheet fracture that generally occurs and propagates in a low temperature range and in a very short time, and it causes significant damage to the entire structure and also has a large effect on the surrounding environment. Even if a crack occurs in a structure due to an unexpected accident or the like, the steel sheet is required to have a characteristic that stops the brittle crack with the steel sheet itself (hereinafter, this characteristic is referred to as a brittle crack propagation stop characteristic).

  A typical steel sheet excellent in brittle crack propagation static characteristics is 9% Ni steel. This steel is a steel plate that is applied to the outer plate of tanks for liquefied natural gas having a boiling point of −162 ° C., and its high Ni content dramatically increases the toughness of the steel plate, resulting in brittle crack propagation stopping properties. Equipped.

  Since the steel is premised on use under a special temperature condition of extremely low temperature, it is acceptable even if there is a significant increase in cost due to containing a high amount of expensive Ni, but as much as liquefied natural gas Application to various large steel structures such as ships, marine structures, construction / civil engineering structures and storage tanks that are not used in a cryogenic environment is difficult from the viewpoint of economy.

  Therefore, a means for improving the brittle crack propagation stop characteristic has been proposed for a general steel sheet for steel structures whose operating temperature does not reach a very low temperature range without increasing the alloy cost. For example, Patent Document 1 focuses on the fact that shear lip (plastic deformation region) generated in a steel surface layer portion when brittle crack propagates is effective in improving the brittle crack propagation stop property, and crystal grains of the shear lip portion are changed. A method is disclosed in which the propagation energy possessed by the brittle cracks that are propagated by miniaturization is absorbed.

During steel plate production, the process of cooling the surface layer part below the Ar ₃ transformation point by controlled cooling after hot rolling, and then stopping the controlled cooling and returning the surface layer part above the transformation point is repeated once or more. In the meantime, by rolling down the steel material, it is repeatedly transformed or processed and recrystallized, and the microstructure of the surface layer portion is made ultrafine to develop a predetermined brittle crack propagation stop characteristic.

  In Patent Document 2, as a method for improving the brittle crack propagation stop property in a steel material mainly composed of ferrite-pearlite, light rolling is performed with a maximum rolling reduction per pass of 12% or less during finish rolling. In order to suppress local recrystallization phenomenon, the ferrite grain size is made as uniform as possible, and the surface portion of the front and back surfaces of the steel material has a ferrite grain having a ferrite equivalent grain diameter of 5 μm or less and an aspect ratio of 2 or more. Has been disclosed.

  However, all the manufacturing methods described in Patent Documents 1 and 2 above have a microstructure effective for brittle crack propagation stopping characteristics by cooling only the steel surface layer part and then recovering the heat and processing during the reheating. Therefore, it is considered that temperature control and management are extremely difficult in the actual manufacturing process of the thick steel plate, and that it is difficult to ensure its homogeneity.

  On the other hand, a method for improving the brittle crack propagation stop property by microstructure control from a viewpoint different from the refinement of the steel sheet surface structure is also known. Specifically, in the controlled rolling process at the time of steel plate production, a texture is developed by applying a reduction to the partially transformed ferrite, thereby causing separation on the fracture surface of the steel plate in a direction parallel to the plate thickness direction. This is a technique for increasing the resistance to brittle fracture by relaxing the stress at the brittle crack tip.

  For example, in Patent Document 3, by controlling rolling, the (110) plane X-ray intensity ratio is set to 2 or more, and the area ratio of coarse particles having an equivalent circle diameter of 20 μm or more is set to 10% or less, thereby exhibiting brittle fracture resistance. An improved steel sheet is described. However, this technique is particularly thick because it requires a corresponding rolling reduction in addition to the concern that the rolling efficiency may be significantly reduced due to the necessity of controlled rolling in a relatively low temperature region after ferrite is formed to some extent. There is a problem that it is difficult to apply to meat.

  By the way, in addition to the brittle fracture propagation stop characteristic, a large heat input welding characteristic is mentioned as an incidental characteristic to the steel plate which is increasing in demand in recent years. Generally, a steel structure obtains a predetermined structure by welding, but as the applied steel sheet becomes thicker, as a measure to increase its construction efficiency, the amount of heat input in one welding (hereinafter, There are an increasing number of cases where high heat input welding (which is referred to as welding heat input) is applied.

  Large heat input welding is performed by high-efficiency methods such as submerged arc welding, electrogas welding, and electroslag welding. Especially when the welding heat input is large, the welding heat affected zone of the steel sheet (hereinafter referred to as the heat affected zone is called HAZ). In some cases, the crystal grains become coarse, or a low toughness and high hardenability structure may be formed, and the HAZ toughness may be remarkably lowered.

  In order to improve the brittle fracture propagation stop characteristics of steel sheets, various controlled rolling is applied to the steel sheets, and the refinement of crystal grains or the promotion of texture development is an extremely effective means. In order to ensure its toughness, the grain refinement effect by the various controlled rolling / cooling processes described above disappears in the vicinity of the melt line portion of the formed HAZ (sometimes referred to as high heat input welding HAZ). Therefore, it is necessary to secure low temperature toughness based on chemical component adjustment, which is not dependent on the manufacturing process.

  As a widely known measure, a technique for suppressing the coarsening of austenite grains by finely dispersing TiN that is relatively stable in a high temperature region during welding, or a Ti oxidation that is stable at high temperatures described in Patent Document 4 There are techniques to disperse things.

  However, the austenite grain refining effect of TiN alone is that TiN dissolves in the vicinity of the melting line portion close to the melting point of the steel, so that it may lead to an increase in solid solution N and significantly reduce low temperature toughness. The technique using oxides has a problem that it is difficult to finely disperse a predetermined oxide uniformly in a steel plate.

  Patent Document 5 proposes to improve the HAZ toughness by fixing solute N as nitride during cooling during welding by using B, which is a nitride-forming element, with a high diffusion rate in a high temperature range. However, when the nitride is already formed at the manufacturing stage of the base material and the solid solution B necessary for improving the HAZ toughness cannot be secured, or because the ferrite formation of the B nitride promotes the formation of ferrite, It has also been pointed out that the strengthening utilized is not sufficient and it is difficult to secure a predetermined strength.

Japanese Patent Publication No. 7-100814 JP 2002-256375 A Japanese Patent Laid-Open No. 10-88280 JP-A-57-051233 JP 2005-2476 A

  Therefore, in order to improve both the brittle crack propagation stopping characteristics and the HAZ toughness in the vicinity of the fusion line part of the high heat input weld HAZ in the thick steel plate, the optimization of the chemical composition for the former and the hot in the steel plate manufacturing stage Both optimization of rolling conditions and measures based on chemical components for the latter are required. In actual machine manufacturing, it is a challenge to reconcile these measures under the premise of high manufacturability on the premise of mass production and high economic efficiency without relying on expensive alloy addition.

  Accordingly, the object of the present invention is to provide a method for producing a thick high-strength steel plate that solves the above-described problems, ensures low temperature toughness of high heat input welding HAZ, and has excellent brittle crack propagation stopping characteristics. And

  In order to solve the above-mentioned problems, the inventors of the present invention have intensively studied by paying attention to the difference between the thermal history during controlled rolling / controlled cooling during steel plate production and the thermal history imparted to the steel plate during high heat input welding. As a result, in the steel plate manufacturing process, the austenite recrystallization temperature range (hereinafter also simply referred to as the recrystallization temperature range) is first rolled on the steel with a component design that is optimized as a measure against high heat input HAZ of high-strength thick steel plates. After that, accelerated cooling to the austenite non-recrystallization temperature range (hereinafter also referred to simply as the non-recrystallization temperature range), followed by non-recrystallization zone rolling, followed by accelerated cooling again results in high strength and excellent low-temperature toughness. It was found that it was obtained.

  That is, in the present invention, as a countermeasure for high heat input welding HAZ toughness, in addition to Ca-based inclusions, B nitride is actively precipitated in the heat input heat history of the high heat input, and ferrite nucleates starting from the precipitate. To ensure toughness by reducing the HAZ high hardenability fraction.

  When these effects are manifested during steel sheet production, the ferrite fraction increases before the controlled rolling is performed, making it difficult to achieve high strength of the steel sheet, and also to ensure brittle crack propagation stopping characteristics in the ferrite main structure In addition, it is necessary to reduce the controlled rolling temperature at a low temperature, and there is a concern that the productivity is significantly reduced.

For this reason, at the time of steel plate production, the temperature range from the rolling end temperature to the non-recrystallization temperature range in the recrystallization temperature range, which is the precipitation temperature range of B nitride, which is the starting point of ferrite nucleation, that is, recrystallization Precipitation of B nitride is prevented as much as possible by accelerated cooling (primary accelerated cooling) of the remaining recrystallization temperature range after the end of rolling in the temperature range, and is carried out after the end of the subsequent non-recrystallization range rolling. By forming a bainite main structure in the second stage accelerated cooling process (secondary accelerated cooling) and developing a transformation texture, it is possible to provide high strength and excellent brittle crack propagation stopping characteristics. I found out. The gist of the present invention is as follows.
1. In mass%, C: 0.030 to 0.080%, Si: 0.01 to 0.10%, Mn: 1.20 to 2.50%, P: 0.008% or less, S: 0.0005 -0.0040%, Al: 0.005-0.1%, Nb: 0.003-0.04%, Ti: 0.003-0.04%, N: 0.003-0.010%, B: 0.0003 to 0.0030%, Ca: 0.0005 to 0.0030%, carbon equivalent Ceq represented by the following formula (1) is 0.33 to 0.45, Ca, S in steel, And O satisfies the following formula (2), and Ti, B, and N in the steel satisfy the following formula (3), and the steel material composed of the remaining Fe and unavoidable impurities is heated to 1000 ° C. or more to obtain austenite After the rolling in the recrystallization temperature range, the temperature range from the start of rolling in the austenite non-recrystallization temperature range to the same Accelerated cooling, subsequently after performing rolling cumulative reduction of 40% or more in the austenite non-recrystallization temperature region, brittle crack propagation comprising a step of secondary accelerated cooling to a temperature range below 500 ℃ from Ar ₃ transformation point or more A method of manufacturing a steel plate for high heat input welding with excellent stopping characteristics.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (2)
−15 <(N—Ti / 3.42-1.269 × B) × 10 ⁴ <15 (3)
However, the element symbols in the above formulas (1) to (3) indicate the content (% by mass) of each element.
2. In addition to the component composition, by mass%, Cu: 1.0% or less, Ni: 1.5% or less, Cr: 1.0% or less, Mo: 0.5% or less, and V: 0.1% or less The manufacturing method of the steel plate for high heat input welding excellent in the brittle crack propagation stop characteristic of 1 characterized by containing 1 type or 2 types or more.
3. The component composition further includes one or more of Mg: 0.0005 to 0.005%, Zr: 0.001 to 0.02%, and REM: 0.001 to 0.02% by mass%. The manufacturing method of the steel plate for high heat input welding excellent in the brittle crack propagation stop characteristic of 1 or 2 characterized by containing.
It has excellent brittle crack propagation stop characteristics as described in any one of 1 to 3, further comprising a step of tempering to a temperature range below the Ac ₁ transformation point after accelerated cooling to a temperature range below 4.500 ° C. Manufacturing method of steel plate for high heat input welding.

  According to the present invention, in addition to optimization of chemical components for securing high heat input welding HAZ toughness, austenite recrystallization temperature range rolling and accelerated cooling to the first stage non-recrystallization temperature range ( (Primary accelerated cooling), followed by non-recrystallization temperature range rolling with a cumulative rolling reduction of 40% or more, and then second stage accelerated cooling (secondary accelerated cooling) to a temperature range of 500 ° C. or lower. Thus, a thick steel plate having both excellent high heat input welding HAZ low temperature toughness and excellent brittle crack propagation stopping characteristics in the steel plate base material portion is obtained, which is extremely useful industrially.

Hereinafter, the present invention will be specifically described. In addition, all% in a chemical component shall be the mass%.
C: 0.030 to 0.080%
C is an element that increases the strength of the steel material, and 0.030% or more of addition is necessary in order to ensure the strength necessary for structural steel. On the other hand, if it exceeds 0.080%, island martensite is likely to be generated in the high heat input welding HAZ, so the upper limit is made 0.080%. Preferably, it is 0.04 to 0.07% of range.

Si: 0.01-0.10%
Si is an element added as a deoxidizer when melting steel, and it is necessary to add 0.01% or more. However, if it exceeds 0.10%, island martensite is generated in the high heat input welding HAZ, and the toughness is liable to be lowered. Therefore, Si is taken as 0.01 to 0.10% of range.

Mn: 1.20 to 2.50%
Like M, Mn is an element that enhances the strength of the steel sheet base metal and is inexpensive compared to other alloy components. Therefore, positive addition of 1.20% or more is effective. If it exceeds 50%, the hardenability becomes excessive, the base metal toughness is lowered, and the weldability is impaired. Therefore, the Mn content is 1.20 to 2.50%. Preferably it is 1.5 to 2.2% of range.

P: 0.008% or less P is one of the elements contained in steel as an impurity. However, in order to reduce the toughness of the steel plate base metal and the high heat input welding HAZ, the economics at the time of material melting are reduced. It is preferable to reduce as much as possible in consideration. For this reason, the amount of P is made into 0.008% or less.

S: 0.0005 to 0.0040%
Like P, S is one of the elements contained in steel as an impurity, but unlike P, when it exists as sulfides such as MnS, CaS, and REM-S, it becomes a nucleus of ferrite formation, resulting in a large heat input. Improve toughness of welded HAZ. This effect can be obtained with a content of 0.0005% or more. On the other hand, an excessive content leads to a large amount of sulfide formation and causes a decrease in base material toughness. Therefore, the S amount is in the range of 0.0005 to 0.0040%.

Al: 0.005 to 0.1%
Al is an element added for deoxidation of steel, and it is necessary to contain 0.005% or more. On the other hand, when it exceeds 0.1%, the amount of inclusions becomes excessive and the toughness of the base material is lowered. Therefore, Al is made 0.005 to 0.1% of range. Preferably, the content is 0.01 to 0.06%.

Nb: 0.003 to 0.04%
Nb has an effect of expanding the non-recrystallization temperature region by addition, and is an element effective for ensuring the strength toughness of the steel plate base material. However, if the addition is less than 0.003%, the above effect is small. On the other hand, if the addition exceeds 0.04%, island martensite is generated in the high heat input welding HAZ, and the toughness is lowered. For this reason, it is set as 0.003 to 0.04% of range. Preferably, it is 0.005 to 0.025% of range.

Ti: 0.003-0.04%
Ti precipitates as TiN during solidification, and is an extremely useful element for increasing the toughness of high heat input weld HAZ, particularly suppressing the coarsening of austenite grains in the weld heat affected zone and becoming a transformation nucleus of ferrite. is there. In order to obtain this effect, 0.003% or more must be added. On the other hand, if added over 0.04%, the precipitated TiN becomes coarse, and the above effect is hardly obtained. Therefore, it is set as 0.003 to 0.04% of range. Preferably, it is 0.005 to 0.025% of range.

N: 0.003-0.010%
N is an element necessary for the formation of TiN described above and the formation of B nitride described later, and is one of the most important elements in the present invention. In order to generate these nitrides in high heat input welding HAZ and effectively contribute to the improvement of toughness, it is necessary to contain 0.003% or more. On the other hand, if the content exceeds 0.010%, depending on the welding heat input conditions, the amount of solute N in the region where TiN dissolves may increase, and on the contrary, the toughness of the welded portion may be reduced. Therefore, the range is 0.003 to 0.010%. Preferably, it is 0.004 to 0.007% of range.

B: 0.0003 to 0.0030%
When B exists in a solid solution state, it is unevenly distributed at the grain boundary to ensure hardenability and contribute to ensuring the strength of the base material. When B exists as a B nitride, it acts as a ferrite nucleus, and high heat input welding. It is one of the most important elements in the present invention that increases HAZ toughness. If the content is less than 0.0003%, the former effect cannot be obtained, and if it exceeds 0.0030%, a large amount of solid solution B exceeding B nitride exists, and conversely, high heat input welding HAZ toughness Causes a drop. Therefore, B is in the range of 0.0003 to 0.0030%.

Ca: 0.0005 to 0.0030%
Ca is an element having an effect of improving the high heat input welding HAZ toughness by fixing S. In order to exert such an effect, it is necessary to contain at least 0.0005%, but even if it exceeds 0.0030%, the effect is saturated, so 0.0005% to 0.0030% And

Ceq: 0.33-0.45%
The steel material according to the present invention contains Ceq (= C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15, in addition to the above-mentioned components satisfying the above composition range, but the elements in the formula The symbol indicates the content (mass%) of each element.) Must be contained in a range of 0.33 to 0.45%.

  If Ceq is less than 0.33%, the required base material strength cannot be obtained even when the controlled rolling / accelerated cooling conditions are adjusted. On the other hand, when Ceq exceeds 0.45%, the amount of island martensite generated in the high heat input welding HAZ becomes extremely large, causing a decrease in toughness. For this reason, it is specified to be 0.33 to 0.45%. Moreover, it is preferably in the range of 0.35 to 0.42%.

0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 However, the element symbol in the formula indicates the content (% by mass) of each element.
In order to improve the high heat input welding HAZ toughness, this formula gives an incidental condition for nucleating ferrite in the HAZ, and controls the balance of Ca, S, and O contents in the steel. This is one of the most important control factors in the present invention, as well as the conditions for securing the ferrite nuclei by balance control of Ti, B, and N amounts described later.

  In order to nucleate ferrite in HAZ, it is necessary to generate as many precipitates as possible with high nucleation ability such as MnS, TiN, and BN. Among these precipitates having a high nucleation ability, MnS forms a Mn dilute band around it to promote ferrite transformation, but usually MnS tends to agglomerate and form, so it is finely dispersed. In addition, inclusions for preferentially precipitating MnS are required.

  In the present invention, attention has been paid to Ca-based inclusions as inclusions that promote precipitation. Ca-based inclusions have a high generation temperature and are crystallized in the solidification stage when the steel is melted. MnS preferentially precipitates at the crystallized phase interface, and aggregation coarsening is suppressed. When (Ca− (0.18 + 130 × Ca) × O) /1.25/S defined by this formula is 0 or less, CaS does not crystallize, S precipitates in the form of MnS alone, and welding The composite sulfide cannot be finely dispersed in the heat affected zone.

  On the other hand, when the value of (Ca− (0.18 + 130 × Ca) × O) /1.25/S is 1 or more, S is completely fixed by Ca, and MnS acting as a ferrite nuclei is CaS. Since it does not precipitate on the surface, the composite sulfide cannot be finely dispersed in the weld heat affected zone. In addition, O is contained in steel as an unavoidable impurity and reduces cleanliness. For this reason, it is desirable to reduce as much as possible in the present invention. In particular, if the O content exceeds 0.003%, CaO-based inclusions become coarse and lower the base material toughness, so the content is preferably made 0.003% or less.

  Further, in the present invention, in order to crystallize Ca as CaS, it is necessary to reduce the amount of O having a strong binding force with Ca before addition of Ca. It is preferable that it is 003% or less. As a method for reducing the amount of residual oxygen, a method such as enhancing degassing or introducing a deoxidizer can be employed.

−15 <(N—Ti / 3.42-1.269 × B) × 10 ⁴ <15 However, each element symbol in the formula represents the content (% by mass) of each element.
This equation is a factor for ensuring high heat input welding HAZ toughness, and is one of the most important control factors in the present invention. As described above, TiN and BN have effective ferrite nucleation ability, but all are nitrides, and if the three elements of Ti, B, and N exist in a solid solution state, the toughness is reduced instead. Therefore, it is necessary to maintain not only the individual content regulations but also the balance.

  This formula is an integer value obtained by subtracting the amount of nitrogen consumed for TiN precipitation and the amount of nitrogen consumed for BN precipitation from the amount of nitrogen contained in steel, and the cooling rate of high heat input welding is Since it is extremely slow, it is assumed that precipitation occurs in a thermal equilibrium according to the stoichiometric ratio.

As a result of studying the toughness of high heat input welding HAZ of steel sheets with various changes in the Ti / B / N ratio, the present inventors have found that (N-Ti / 3.42-1.269 × B ) When the value of × 10 ⁴ is −15 or less, Ti or B is excessive, and in the former, precipitation of TiC is manifested, and in the latter, an increase in hardenability due to solute B is manifested, thereby reducing HAZ toughness. When the value of (N-Ti / 3.42-1.269 × B) × 10 ⁴ is 15 or more, conversely, less nitrogen is consumed for precipitation, and excess nitrogen exists in the HAZ as a solid solution. As a result, the toughness was lowered in spite of the ferrite main structure. Therefore, the specified range of this formula is specified as described above.

  Although the basic component composition of the present invention is as described above, in order to further improve desired properties, one or more of Cu, Ni, Cr, Mo, V, Mg, Zr, and REM are contained as selective elements. be able to.

Cu: 1.0% or less Cu is an element that can be contained in order to increase the strength, but if contained in excess of 1.0%, the quality of the surface of the steel sheet base material is deteriorated due to hot brittleness. When it contains, the quantity shall be 1.0% or less of range.

Ni: 1.5% or less Ni is an element that can improve the toughness while increasing the strength of the base material. When the content exceeds 1.5%, the effect is saturated and economically disadvantageous. Therefore, when it is contained, the amount is in the range of 1.5% or less, preferably in the range of 1.0% or less. To do.

Cr: 1.0% or less Cr is an effective element for increasing the strength. However, if it exceeds 1.0%, the toughness of the base metal is deteriorated. % Or less.

Mo: 0.5% or less Mo is an element effective for increasing the strength of the base metal. However, when it is contained in excess of 0.5%, the toughness is significantly deteriorated and the economy is impaired. The amount is in the range of 0.5% or less.

V: 0.1% or less V is an element effective for increasing the strength of the base metal. However, if the content exceeds 0.1%, the toughness is remarkably deteriorated. The range is 1% or less.

Mg: 0.0005-0.005%, Zr: 0.001-0.02% and REM: 0.001-0.02%
Mg, Zr, and REM have the effect of fixing S in the steel to improve the toughness of the steel sheet. Mg, which is a relatively strong sulfide-forming element, is 0.0005% or more, and 0 for Zr and REM. Each content is effective when the content is 0.001% or more. However, when each content exceeds 0.005%, 0.02%, and 0.02%, the amount of inclusions in the steel increases and the toughness is deteriorated instead. Therefore, when these elements are contained, Mg is in the range of 0.0005 to 0.005%, Zr is in the range of 0.001 to 0.02%, and REM is in the range of 0.001 to 0.02%. The balance other than the above components is composed of Fe and inevitable impurities.

  In the production of the steel according to the present invention, a slab for producing a steel sheet is produced by using a conventional process such as a continuous casting method or an ingot forming method, and a molten steel melted by a conventional melting method such as a converter or an electric furnace. It is preferable to use a material. Hereinafter, the reason for limiting the steel sheet manufacturing conditions will be described. The steel material temperature in the present invention is the average temperature of the steel material surface and the central part (1/2 part of the plate thickness).

Heating temperature: 1000 ° C. or higher After casting, the slab after casting is cooled to room temperature or, if possible, kept in a high temperature state and charged in a heating furnace, and the heating temperature is regulated to 1000 ° C. or higher. . In the heating of the slab, Nb carbonitride was mainly dissolved, and the lower limit temperature was set to 1000 ° C. from the viewpoint of sufficiently securing the solid solution Nb. Further, although the upper limit side of the heating temperature is not specified, when it is excessively high, austenite grains become coarse during heating and adversely affect the toughness of the base metal. Therefore, it is usually 1250 ° C. or lower, preferably 1200 ° C. or lower. .

Rolling in the austenite recrystallization temperature range Rolling in the austenite recrystallization temperature range is necessary for refining the austenite grains at the time of heating, and it is desirable to carry out at least 1 pass, preferably 20% or more.
In addition, it is desirable to perform the rolling at a low temperature side of the austenite recrystallization temperature range if possible.

The primary accelerated cooling of the temperature range from the end of rolling in the austenite recrystallization temperature range to the start of rolling in the austenite non-recrystallization temperature range is one of the most important items in the present invention. As described above, in the present invention, in order to improve the toughness of the high heat input welding HAZ, the grain size refinement by TiN and the ferrite transformation promoting effect by the formation of B nitride are utilized. In the process, particularly when a large amount of B nitride is generated, the solid solution B for securing the hardenability of the steel plate disappears, while ferrite nucleation is likely to occur from the precipitated B nitride. There is a possibility that the ferrite fraction occupying the rolled structure increases and a predetermined strength cannot be obtained.

Therefore, the cooling rate in the temperature range from the recrystallization temperature range to the non-recrystallization temperature range, which is the temperature range in which B nitride is generated, in the cooling process at the time of steel plate production is increased as much as possible. is necessary. Normally, this process is air-cooled as the temperature reduction waiting time of hot rolling, but in the present invention, the primary accelerated cooling is performed after the rolling in the austenite recrystallization temperature range, whereby the non-recrystallization temperature range which is the next step. Let's move to the rolling process in a short time.
In this primary accelerated cooling, it is desirable to control the lower limit of the cooling rate so that B nitride does not precipitate during the cooling. The inventors found that the critical cooling rate for precipitation of B nitride is approximately 2 ° C./second within the range of Ti, B, and N described above, and if the cooling is performed at a cooling rate equal to or higher than the critical cooling rate, non-recrystallization occurs. It was experimentally found that precipitation of B nitride was suppressed until zone rolling, and a solid solution B amount sufficient to ensure hardenability during secondary accelerated cooling was obtained. Therefore, in the primary accelerated cooling, a cooling rate of 2 ° C./second or more is preferable.

  In that case, a cooling rate faster than air cooling can be achieved, for example, by an accelerated cooling facility by water cooling, or a so-called deske facility that removes oxide scale generated on the surface of the steel sheet during rolling, and 2 ° C./second or more. The cooling rate can be as follows.

Rolling at a cumulative reduction ratio of 40% or more in the austenite non-recrystallization temperature region Following the primary accelerated cooling, rolling is performed in the austenite non-recrystallization temperature region. In this rolling, when the rolling reduction is small, a predetermined base material toughness cannot be obtained. For this reason, the lower limit of the cumulative rolling reduction is defined as 40%. Moreover, although the one where a rolling reduction is higher is preferable, about 80% becomes an upper limit industrially. The rolling end temperature is preferably not less than the Ar ₃ transformation point in order to avoid two-phase region rolling.

Secondary rolling cooling after non-recrystallization temperature range rolling to a temperature range from Ar ₃ transformation point to 500 ° C. or less Unlike the above-mentioned primary accelerated cooling for suppressing B nitride formation, secondary accelerated cooling in this step is This is a process for phase transformation of the austenite structure processed by controlled rolling. In order to complete the phase transformation, it is necessary to cool to a temperature range of 500 ° C. or lower, so the upper limit of the cooling temperature was set to 500 ° C. As for the cooling rate of the secondary accelerated cooling, it is desirable to set a lower limit cooling rate for satisfying a predetermined required strength. In order to achieve a suitable strength even in a relatively thick steel plate to which large heat input welding is applied, which is the main object of the present invention, strong cooling of 5 ° C./sec or more is preferable. The cooling method is not particularly limited, but cooling by water cooling is preferable.

Tempering After the secondary accelerated cooling, a tempering treatment can be performed as necessary. Tempering is carried out mainly for the purpose of optimizing the balance between strength and toughness and reducing residual stress on steel that has been quenched by secondary accelerated cooling, and when it is carried out at a temperature below the Ac ₁ transformation point. Do.

The Ar ₃ and Ac ₁ transformation points vary depending on the steel components. The Ar ₃ and Ac ₁ transformation points can be obtained by the following equation. However, in each formula, each element symbol indicates the content (% by mass) of each element.
Ar ₃ = 910-273C-74Mn-56Ni-16Cr-9Mo-5Cu
_{Ac 1 = 751-26.6C + 17.6Si-11.6Mn} -169Al-23Cu-23Ni + 24.1Cr + 22.5Mo + 233Nb-39.7V-5.7Ti-895B
On the other hand, the lower limit temperature of the austenite recrystallization temperature region is influenced by the crystal grain size, processing history, strain amount, etc. in addition to the steel composition, but is generally in the range of 800 to 950 ° C. By conducting a preliminary test and investigating in advance, the lower limit temperature can be estimated. Hereinafter, the effects of the present invention will be described in detail with reference to examples.

  Steels having the compositions shown in Table 1 were melted in a converter and then made into slabs (steel materials) by a continuous casting method, and steel sheets having a thickness of 40 to 80 mm were produced according to the controlled rolling and controlled cooling conditions shown in Table 2. Note that the rolling in the austenite recrystallization temperature range was performed under the conditions of one pass or more and a cumulative reduction ratio of 20% or more. It was carried out by temporary accelerated cooling water cooling. Secondary accelerated cooling was performed at a cooling rate of 5 ° C./sec or more by water cooling.

In Table 1, steel numbers 1 to 10 are steels within the scope of the present invention, and steel numbers 11 to 20 have a component composition and / or any of formulas (1) to (3) outside the scope of the present invention. It is steel. Furthermore, in Table 2, the branch numbers following the steel numbers start with A (A1, A2) are due to the controlled rolling / cooling conditions according to the present invention, and the branch numbers start with B (B1, B2). Is a comparative example in which any of the manufacturing conditions falls outside the scope of the present invention.

  Among the steel numbers 1 to 10, the steel plate whose branch number starts with B is a comparative example in which the chemical composition is within the scope of the present invention but the production method is outside the scope of the present invention.

About the manufactured thick steel plate, a tensile test piece having a parallel part diameter of 6 mmφ is taken from the position of 1/4 of the plate thickness direction as the longitudinal direction, and is tensioned according to the provisions of JIS Z 2241 (1998). The test was carried out to determine the tensile strength (hereinafter referred to as TS) and 0.2% proof stress (hereinafter referred to as YS). In the present invention, a high strength steel plate is assumed as the object, and the threshold value of the strength target is TS: 520 N / mm ² .

In addition, a Charpy impact test piece having a V-notch standard size was collected from a 1/4 position in the plate thickness direction and in a direction parallel to the rolling direction in accordance with JIS Z 2202 (1998), and JIS Z 2242 (1998). In order to evaluate the brittle crack propagation stop property of the steel sheet, a temperature gradient type ESSO test was conducted to obtain a fracture surface transition temperature (hereinafter referred to as vTrs). 10 ° C.), and the threshold of the crack propagation stop characteristic target was evaluated as 7000 N / mm ^3/2 .

  Furthermore, in order to evaluate the toughness of the high heat input welding HAZ, a test piece of width 80 mm × length 80 mm × thickness 15 mm is taken from each thick steel plate, heated to 1450 ° C., and then cooled to 800 to 500 ° C. in 400 seconds. After the heat treatment was applied, a 2 mm V notch Charpy test piece was collected and subjected to a Charpy impact test in the same manner as described above. The impact test temperature was −40 ° C., and the threshold value of the toughness target was 50 J as an average value of absorbed energy (hereinafter referred to as vE-40 ° C.) of three test pieces at −40 ° C.

  Table 3 shows the steel plate base material characteristics, brittle crack propagation stop characteristics, and high heat input weld HAZ toughness evaluation results. In the steel samples having the steel numbers 1 to 10 and the branch numbers starting with A according to the present invention, good values are obtained for both the brittle crack propagation stop characteristic and the high heat input welding HAZ characteristic of the base material.

  For steel numbers 1 to 10 and branch numbers beginning with B, the chemical composition is within the scope of the present invention, so the high heat input welding HAZ characteristics are satisfied, but the manufacturing conditions are out of the range, so the brittleness of the base metal The crack propagation stop property is inferior, and in steel numbers 11 to 20 (the branch number starts with A or B), the toughness of the high heat input weld HAZ is inferior because the component range is outside the scope of the present invention. It was recognized that

Claims (4)

In mass%, C: 0.030 to 0.080%, Si: 0.01 to 0.10%, Mn: 1.20 to 2.50%, P: 0.008% or less, S: 0.0005 -0.0040%, Al: 0.005-0.1%, Nb: 0.003-0.04%, Ti: 0.003-0.04%, N: 0.003-0.010%, B: 0.0003 to 0.0030%, Ca: 0.0005 to 0.0030%, carbon equivalent Ceq represented by the following formula (1) is 0.33 to 0.45, Ca, S in steel, And O satisfies the following formula (2), and Ti, B, and N in the steel satisfy the following formula (3), and the steel material composed of the remaining Fe and unavoidable impurities is heated to 1000 ° C. or more to obtain austenite After the rolling in the recrystallization temperature range, the temperature range from the start of rolling in the austenite non-recrystallization temperature range to the same Brittle crack having a step of performing secondary accelerated cooling to a temperature range of not less than Ar ₃ transformation point to 500 ° C. or less after performing rolling accelerated cooling and subsequently rolling at a cumulative reduction ratio of 40% or more in the austenite non-recrystallization temperature range A method of manufacturing a steel plate for high heat input welding with excellent propagation stop characteristics.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (2)
−15 <(N—Ti / 3.42-1.269 × B) × 10 ⁴ <15 (3)
However, the element symbols in the above formulas (1) to (3) indicate the content (% by mass) of each element.   In addition to the component composition, by mass%, Cu: 1.0% or less, Ni: 1.5% or less, Cr: 1.0% or less, Mo: 0.5% or less, and V: 0.1% or less The manufacturing method of the steel plate for high heat input welding excellent in the brittle crack propagation stop characteristic of Claim 1 characterized by containing 1 type, or 2 or more types.   The component composition further includes one or more of Mg: 0.0005 to 0.005%, Zr: 0.001 to 0.02%, and REM: 0.001 to 0.02% by mass%. The manufacturing method of the steel plate for high heat input welding excellent in the brittle crack propagation stop characteristic of Claim 1 or 2 characterized by including. It has excellent brittle crack propagation stop characteristics according to any one of claims 1 to 3, further comprising a step of tempering to a temperature range below the Ac ₁ transformation point after accelerated cooling to a temperature range below 500 ° C. Manufacturing method of steel plate for high heat input welding.