KR20080057432A - Thick steel having excellent high temperature ductility for high heat input welding - Google Patents

Abstract

A thick steel sheet for high heat input welding is provided to improve high temperature ductility by controlling the volume of AIN precipitate and enhance toughness of a welding heat effecting part, using fine TiN precipitate, Ti-Nb and Ti-V composite precipitate. A thick steel sheet having excellent high temperature ductility for high heat input welding is composed of, by weight %, C:0.03-0.17%, Si:0.01-0.5%, Mn:0.4-2.0%, Ti:0.005-0.2%, Ai:0.005-0.05%, N:0.005-0.020%, B:0.0003-0.01%, W:0.001-0.2%, Nb:0.005-0.1%, V:0.005-0.1%, P: less than 0.03%, S: less than 0.03%, O: less than 0.005%, and the remainder Fe and unavoidable impurities. Ti, N, B, Nb, V, and Al satisfy the followings, 1.2<=Ti/N<=4, 3<=N/B<=40, 0.3<=Nb/N<=9, 0.5<=V/N<=7, and 8x10^-5<=Al*N<=5x10^-4. A fine tissue of a base is composed of ferrite and bainite.

Description

Thick steel having excellent high temperature ductility for high heat input welding}

Japanese Patent Laid-Open No. 11-140582

Japanese Patent Laid-Open No. 10-298708

Japanese Patent Laid-Open No. 9-194990

The present invention relates to structural steels mainly used in welding structures, such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes. More specifically, the high temperature ductility is improved by controlling the amount of AlN precipitates, and fine TiN precipitates and Ti-Nb and Ti-V composite precipitates are used for the plate steel for high heat input welding, which has improved the toughness of the weld heat affected zone. will be.

Recently, steels used in accordance with the trend of high-rise buildings and structures have been replaced by thick steels. In order to weld such a thick material, high efficiency welding is inevitable. As a technique for welding a thickened steel material, a high heat input thermally-merged welding method and an electro-welding method capable of one-pass welding are widely used. In addition, in the field of shipbuilding and bridges, even in the case of welding a steel plate having a plate thickness of 25 mm or more, the above-described high heat input welding method capable of 1 pass welding is applied to improve welding productivity.

In general, in welding, the larger the amount of heat input, the larger the amount of welding, so that the number of welding passes decreases. Therefore, it is advantageous to enable high heat input welding in consideration of welding productivity. In other words, increasing the amount of heat input in the welding will be able to widen the range of use. The range of high heat input currently in use is about 100-200kJ / cm, but in order to weld steel with more than 50mm thicker plate thickness, it is possible to have a super heat input range of 200-800kJ / cm.

 Heat Affected Zone in which the heat input of steel is formed during welding, particularly the weld heat affected zone near the fusion boundary, is heated to a temperature close to the melting point by the heat input of the weld. Accordingly, since the grains of the weld heat affected zone grow and coarse, and microstructures that are vulnerable to toughness such as upper bainite and martensite are formed during the cooling process, the weld heat affected zone is the site where the toughness of the weld deteriorates most.

 Therefore, in order to secure the stability of the welded structure, it is necessary to suppress the austenite grain growth of the weld heat affected zone and keep it fine. As a means to solve this problem, there is known a technique for delaying grain growth of the weld heat affected zone during welding by appropriately dispersing an oxide or Ti-based carbonitride, which is stable at a high temperature, into a steel material, and the prior art has disclosed a Japanese patent. Publication Nos. 11-140582, 10-298708, and 9-194990.

The Japanese Laid-Open Patent Publication No. 11-140582 is a representative technique using TiN precipitates, and the impact toughness at 200 ° C. is about 200J when the heat input amount of 100 J / cm (maximum heating temperature 1400 ° C.) is applied (a base material is about 300 J). Phosphorus structural steels are disclosed. In the prior art, Ti / N was substantially managed at 4-12 so that TiN precipitates of 0.05 μm or less were 5.8 × 10 ³ to 8.1 × 10 ⁴ pieces / mm 2, and TiN precipitates of 0.03 to 0.2 μm were 3.9 × 10 ³ pieces. The toughness of the welded portion is secured by making the ferrite fine by precipitating at / mm2 to 6.2x10 ⁴ pieces / mm2.

 However, according to the prior art, when 100 kJ / cm high heat input welding is applied, the toughness of the base material and the heat affected zone is generally low (the base material: 320J, the heat affected zone: 220J with the highest impact toughness of 0 ° C.), and the base material. Since the toughness difference between the and the heat affected zone is about 100J, there is a limit in securing the reliability of the steel structure due to superheated welding of the thickened steel. In addition, as a method for securing the desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or higher and quenched and then reheated for hot rolling. do.

 Japanese Patent Application Laid-Open No. 10-298708 discloses a technique using MgO-TiN composite precipitates, but the toughness of the impact resistance of the welding heat-affected zone 0 ° C. is 130J when the high heat input welding of about 100 kJ / cm is applied. In addition, Japanese Patent Application Laid-Open No. 9-194990 manages the ratio of Al and O to 0.3 ≦ Al / O ≦ 1.5 in low nitrogen steel (N = 0.005%), and uses a composite oxide of Al, Mn, and Si. However, when the high heat input welding of about 100 kJ / cm is applied, the weld heat affected zone transition temperature is -50 ° C, which is not good at toughness.

 Until now, there are many known techniques for improving the toughness of the weld heat affected zone during high heat input welding using TiN precipitates and Al-based or MgO oxides. However, the toughness of the weld heat affected zone during super long heat welds maintained at a high temperature of 1350 ° C. or more for a long time is known. No significant improvements have been made yet. In particular, there is little technology for improving the high temperature ductility of the high heat input welding steel containing a large amount of precipitates. Therefore, if the above problems can be solved, it will be possible to significantly improve the product quality of the high heat input welding steel.

The present invention is to improve the conventional problems described above, by controlling the amount of AlN precipitates to improve the high temperature ductility, by using the fine TiN precipitates and Ti-Nb and Ti-V composite precipitates, welding heat effect by the tissue control It is an object of the present invention to provide a thick plate steel for high heat input welding with improved negative toughness.

The present invention for achieving the above object, in the weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.005-0.05%, N: 0.005-0.020%, B: 0.0003-0.01%, W: 0.001-0.2%, Nb: 0.005-0.1%, V: 0.005-0.1%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less It is composed of the remaining Fe and other unavoidable impurities, including, Ti, N, B, Nb, V and Al are 1.2≤Ti / N≤4, 3≤N / B≤40, 0.3≤Nb / N≤9, Substituting excellent high temperature ductility, which satisfies 0.5≤V / N≤7, 8 × 10 ⁻⁵ ≦ Al * N ≦ ⁵ × 10 ⁻⁴ , and the microstructure of the base material is composed of a composite structure of ferrite and bainite of 20 μm or less. It relates to a thick plate steel for heat welding.

EMBODIMENT OF THE INVENTION Hereinafter, this invention is demonstrated in detail.

 In the present invention, the term "prior austenite" refers to austenite formed in the weld heat affected zone when the heat input welding is applied to the steel (base material), and is formed in the manufacturing process of the steel (hot rolling process). It is used for convenience to distinguish it from being austenite.

The inventors have found that excessive formation of AlN precipitates dispersed in the base material of the steel reduces the high temperature ductility of the steel, thereby causing a large amount of defects on the surface of the steel during the production of the steel. Therefore, in the present invention, by limiting the amount of AlN that can cause defects on the steel product surface to improve the high temperature ductility of the steel, it is possible to minimize the defect rate of the product by dramatically reducing the defects on the steel surface.

In addition, the inventors of the present invention have found that the toughness of the weld heat-affected zone changes based on the grain size (about 80 μm) of the critical austenite. .

  In this regard, the present inventors have been able to derive the following method for securing the toughness of the weld heat affected zone.

[1] TiN precipitates and Ti-Nb and Ti-V composite precipitates are used.

[2] With the initial ferrite grain size of the steel below the critical level, when the high heat input welding is applied, the former austenite of the heat affected zone is refined to 80 μm or less.

[3] Enhancement of ferrite formation in the austenite mouth of weld heat affected zone by using Ti-Nb and Ti-V complex precipitates and BN precipitates, and promote the transformation of polygonal or acicular ferrites, especially in the austenite Increase the toughness improvement effect.

[4] Ti-Nb and Ti-V complex precipitates and BN precipitates are used to reduce the content of soluble nitrogen during cooling of the weld heat affected zone.

[1], [2], [3], and [4] will be described in more detail below.

[1] TiN precipitate and Ti-Nb and Ti-V composite precipitate management

When heat input welding is applied to structural steels, the weld heat affected zone near the melting line is heated to a high temperature of about 1400 ° C or higher, and TiN precipitates deposited in the base metal are partially dissolved by welding heat or Oswald ripening (Ostwald ripening, Small precipitates decompose and diffuse into larger precipitates, causing larger precipitates to become larger). Some precipitates are degraded or some precipitates become coarse, and the number of TiN precipitates is significantly reduced, thereby suppressing the growth of the austenite grains. The effect disappears.

The inventors have found that this phenomenon is caused by the diffusion of TiN precipitates dispersed in the base metal by the dissolution of the dissolved Ti atoms by the heat of welding, and the high temperature stability of the TiN precipitates according to the ratio of Ti / N was investigated. In the low Ti / N ratio), it was found that the dissolved Ti concentration and diffusion rate of the dissolved Ti atoms were reduced, and the high temperature stability of the TiN precipitates was improved. That is, when the ratio of Ti and N (Ti / N) is in the range of 1.2 to 4, the amount of solid solution Ti is extremely reduced and the high temperature stability of the TiN precipitate is greatly improved, so that the fine TiN precipitate having a size of 0.01 to 0.1 μm is 0.5. Important results were obtained that were distributed at 1.0 × 10 ⁷ / mm 2 or more at intervals of less than or equal to μm. Increasing the nitrogen content at the same Ti content causes all Ti atoms to be easily combined with the nitrogen atoms, and in the high nitrogen environment, the amount of solid solution Ti decreases, making TiN precipitates more stable at higher temperatures than in the case of low nitrogen content. It is analyzed that this is because the solubility product is lowered. In view of the fact that the present invention can promote aging due to the presence of solid solution N in a high nitrogen environment, the ratio of N / B, Nb / N, and V / N is collectively managed so that N is BN, Ti-Nb. And composite precipitated with Ti-V.

Furthermore, in the present invention, the Ti-Nb and Ti-V composite precipitates which are stable at high temperature together with the fine and uniformly distributed TiN precipitates are used together. Ti-Nb and Ti-V composite precipitates are dispersed in the base material to suppress the growth of austenite grains in the weld heat affected zone during welding, as well as to promote the polygonal ferrite transformation in the austenite grains during the cooling process. Negative toughness can be improved.

[2] ferrite grain size management

According to the study of the present invention, in order to average the size of the old austenite in the welding heat affected zone to 80㎛, the size of the ferrite in the base material structure of ferrite + bainite and finer to about 20㎛ or less Is important. At this time, the refinement of the ferrite may be obtained by inhibiting the growth of ferrite grains generated during the cooling process using carbides (Fe ₃ C, NbC, VC, WC) together with austenite grain refinement by hot processing during hot rolling. .

[3] microstructure of weld heat affected zone

The facts of the present invention reveal that the toughness of the weld heat affected zone is not only the former austenite grain size when the base material is heated above about 1400 ° C., but also the amount and size of ferrite (more than 70%) precipitated at the old austenite grain boundary. (Less than 20㎛) And its shape has an important effect. In the present invention, the Ti / N ratio is 1.2-4, the Nb / N ratio is 0.3-9, and the V / N ratio is 0.5-7, and the number of TiN, Ti-Nb and Ti-V composite precipitates and precipitates such as BN is It generates a large amount of fine ferrite in the mouth of the old austenite, the ferrite is mostly polygonal (ferrite) ferrite can greatly improve the toughness of the heat affected zone.

[4] free nitrogen content reduction in welded heat affected zones

In fact, in the study of the present invention, a part of the precipitates dispersed in the base material during the high heat input welding is decomposed by the welding heat, so that the solid solution nitrogen content in the weld heat affected zone is increased, thereby reducing the toughness of the weld heat affected zone. Likewise, the formation of TiN precipitates, Ti-Nb and Ti-V complex precipitates in a high nitrogen environment can reduce the solid solution nitrogen content in the weld heat affected zone, thereby suppressing the aging effect caused by solid solution nitrogen, thereby contributing to improving the toughness of the weld heat affected zone. .

Hereinafter, the composition range of the steel component of the present invention will be described.

The content of carbon (C) is preferably 0.03 to 0.17%.

If C is less than 0.03%, it is not preferable because other alloying elements are required to secure strength as structural steel. In addition, when it exceeds 0.17%, microstructures susceptible to toughness such as upper bainite, martensite and degenerate pearlite are transformed during cooling to lower the low temperature impact toughness of structural steel, and also increase the crack susceptibility of welded parts. It is not preferable.

The content of silicon (Si) is preferably 0.01 to 0.5%.

When the Si content is less than 0.01%, the deoxidation effect of the molten steel is insufficient in steelmaking process, and the corrosion resistance of the steel is deteriorated. When the content of Si exceeds 0.5%, the effect is saturated, and the quenchability during cooling after rolling is increased. It is not preferable because it promotes transformation of island martensite to lower the low temperature impact toughness and affect the weld cracking susceptibility.

The content of manganese (Mn) is preferably 0.4 to 2.0%.

Mn is an effective element that improves deoxidation, weldability, hot workability and strength in steel, and precipitates in the form of MnS + VN around Ti-based oxides to produce needle-shaped and polygonal ferrites effective for improving the toughness of the weld heat affected zone. Influencing element. Such Mn forms a solid solution to form a solid solution in the matrix structure to strengthen the matrix solution to secure strength and toughness. However, when the Mn content exceeds 2.0%, tissue heterogeneity due to Mn segregation has a detrimental effect on the toughness of the weld heat affected zone, rather than the effect of solid solution strengthening. In addition, macro segregation and micro segregation may occur according to the segregation mechanism of the steel during solidification, thereby promoting the formation of a central segregation zone in the center of the steel sheet, which may act as a cause of generating low temperature transformation structure in the center of the base metal.

The content of titanium (Ti) is preferably 0.005 to 0.2%.

Ti is an indispensable element in the present invention because it combines with N to form a fine TiN precipitate that is stable at high temperature as well as to form a Ti-Nb composite precipitate. In order to obtain such a fine TiN, Ti-Nb and Ti-V composite precipitate effect, it is preferable to add more than 0.005% of Ti, but if it exceeds 0.2%, coarse TiN precipitate and coarse Ti-Nb and Ti-V complexes in molten steel It is not preferable because the precipitate is formed in the molten steel and mixed in the slab and the base material to prevent the growth of the weld heat-affected zone austenite grains during welding.

The content of aluminum (Al) is preferably 0.005-0.05%.

Al is an element that is necessary as a deoxidizer and forms an Al oxide by reaction with oxygen, thereby preventing Ti from reacting with oxygen, thereby being effective for Ti to form fine TiN precipitates. Therefore, it is good to make Al content into 0.005% or more. However, if the content exceeds 0.05%, a large amount of AlN may be precipitated to lower high-temperature ductility during the manufacturing process of the steel, thereby causing surface defects of the steel.

The content of nitrogen (N) is preferably 0.005-0.02%.

The N is an indispensable element for forming TiN, BN, (Ti-Nb) N and (Ti-V) N, and the like, and the maximum suppression of the growth of the austenite grains in the weld heat-affected zone during the high heat input welding, The amount of precipitates such as AlN, BN, and NbN is increased. In particular, the content of TiN and AlN precipitates and the spacing of precipitates, the distribution of precipitates, the frequency of complex precipitation with oxides, and the high temperature stability of the precipitates themselves are significantly affected. Therefore, the content is preferably set to 0.005% or more. However, if the nitrogen content exceeds 0.02%, the effect is saturated, and toughness decreases due to the increase in the amount of solid solution nitrogen distributed in the weld heat affected zone, and it is mixed in the weld metal due to dilution during welding, which may cause the toughness of the weld metal. Can be.

The content of boron (B) is preferably 0.0003-0.01%.

B is a very effective element for producing polygonal ferrite at grain boundaries as well as acicular ferrite having excellent toughness in grains. B forms a BN precipitate, which hinders the growth of the old austenite grains and forms Fe carbide in the grain boundary and in the mouth to promote ferrite transformation of acicular and polygons having excellent toughness. If the B content is less than 0.0003%, such an effect cannot be expected, and if it exceeds 0.01%, the hardenability increases, which may cause hardening of the weld heat affected zone and low temperature cracking.

The content of tungsten (W) is preferably 0.001-0.2%.

W is an element that uniformly precipitates tungsten carbide (WC) in the base material after hot rolling to effectively suppress ferrite grain growth after ferrite transformation, and also suppress the growth of the initial austenite grains heated during the welding heat affected zone. If the content is less than 0.001%, there is less distribution of tungsten carbide for suppressing ferrite grain growth upon cooling after hot rolling, and the effect is saturated when more than 0.2% is added.

The content of niobium (Nb) is preferably 0.005-0.1%.

The Nb is a useful element that combines with Ti to form a Ti-Nb composite precipitate to promote ferrite transformation in the weld heat affected zone. To form a fine Ti-Nb composite precipitate, Nb content of 0.005% or more is preferably added. Do. However, if the Nb content exceeds 0.1%, the low temperature transformation structure in the weld heat affected zone may increase, which may adversely affect the mechanical properties.

The content of vanadium (V) is preferably 0.005-0.1%.

The V is a useful element that combines with Ti to form a Ti-V composite precipitate to promote ferrite transformation in the austenite grains in the weld heat affected zone, so that a V content of 0.005% or more is added to form a fine Ti-V composite precipitate. It is preferable. However, if the V content exceeds 0.1%, it is not preferable because it increases the low-temperature transformation structure in the weld heat affected zone, which adversely affects the mechanical properties and adversely affects the weld cracking susceptibility.

The content of phosphorus (P) and sulfur (S) is preferably 0.030% or less.

P is preferably managed as low as possible because it is an impurity element that promotes high temperature cracking during welding and central segregation during rolling. In order to improve the toughness of the base metal, the toughness of the weld heat affected zone, and to reduce the center segregation, it is recommended to manage it to 0.03% or less.

S is preferably managed as low as possible because it forms a low melting point compound such as FeS when present in a large amount. In order to reduce the base material toughness, weld heat affected zone toughness and central segregation, it is recommended that the S content be 0.03% or less. Particularly, in the case of sulfur, it precipitates in the form of MnS around Ti-based oxides, which affects the formation of needle-shaped and polygonal ferrites, which are effective for improving the toughness of welding heat affected zones. It is desirable to limit it to 0.03% or less.

The content of oxygen (O) is preferably 0.005% or less.

O is an element that forms Ti oxide by reacting with Ti in molten steel, and this Ti oxide promotes the transformation of acicular ferrite from guustenite in the weld heat affected zone. When the content of oxygen (O) exceeds 0.005%, coarse Ti oxides and other oxides such as FeO are generated, which is not preferable because it affects the weld heat affected zone.

The present invention is composed of Fe and other unavoidable impurities in addition to the above components.

Hereinafter, in the present invention, the relation of Ti, N, B, Nb, V and Al will be described in detail.

It is preferable to limit the ratio of Ti / N to 1.2-4.

In the present invention, the Ti / N ratio is lowered to 4 or less, which has two advantages. First, it is possible to increase the amount of TiN, that is, the number of TiN precipitates. In other words, if the nitrogen content is increased at the same Ti content, all the Ti atoms dissolved in the cooling process during the playing process combine with the nitrogen atom, thereby increasing the fine TiN precipitation. Second, TiN is stable at high temperatures. That is, since the solubility product which shows the stability of the precipitate at high temperature such as the weld heat affected zone becomes smaller, the precipitate such as high nitrogen TiN is more stable than the case where the nitrogen content is low. On the other hand, when the Ti / N ratio is higher than 4, coarse TiN is crystallized in molten steel, which is a steelmaking process, and a uniform distribution of TiN is not obtained. Affects bad toughness. If the Ti / N ratio is less than 1.2, the amount of solid solution nitrogen in the base metal increases, which is detrimental to the toughness of the weld heat-oriented part.

It is preferable to limit the ratio of N / B to 3-40.

In the present invention, when the N / B ratio is less than 3, the amount of precipitation of BN that promotes the ferrite transformation of polygons at the old austenite grain boundary during the cooling process after welding is insufficient, and when the N / B ratio exceeds 40, the effect is saturated. The amount of solid solution nitrogen may increase to reduce the toughness of the weld heat affected zone.

It is preferable to limit Nb / N ratio to 0.3-9.

When the Nb / N ratio is less than 0.3, the number of Ti-Nb complex precipitates required for austenite grain growth inhibition is insufficient, and the Nb ratio contained in the composite precipitates becomes small, resulting in loss of function as a needle-like ferrite nucleation site, thus affecting welding heat. The needle-like ferrite phase fraction effective for improving negative toughness decreases. In addition, when the Nb / N ratio exceeds 9, the austenite grain growth inhibitory effect is saturated, and the ferrite loses its function as nucleation site.

It is preferable to limit V / N ratio to 0.5-7.

If the V / N ratio is less than 0.5, the number of Ti-V complex precipitates required for austenite grain growth inhibition is insufficient, and the V ratio contained in the composite precipitates becomes small, thus losing the function as a needle-like ferrite nucleation site and affecting welding heat. The needle-like ferrite phase fraction effective for improving negative toughness decreases. In addition, when the V / N ratio exceeds 7, the austenite grain growth inhibitory effect is saturated, and the ferrite loses its function as nucleation site.

The Al * N value is preferably limited to 8 × 10 ⁻⁵ to 5 × 10 ⁻⁴ .

When the Al * N value is less than 8 × 10 ^−5, the amount of AlN precipitates that can generate a large amount of fine ferrite in the old austenite grain is greatly reduced. On the other hand, when it exceeds 5 × 10 ⁻⁴ , the amount of AlN precipitates is excessively large, which greatly reduces the high temperature ductility of the steel, and may cause a large amount of defects on the surface of the steel product.

In the present invention, at least one of Ni, Cu, Mo, and Cr may be further added to the steel formed as described above to further improve the mechanical properties.

The content of nickel (Ni) is preferably 0.1 to 3.0%.

Ni is an effective element which improves the strength and toughness of the base material by solid solution strengthening. In order to achieve this effect, the Ni content is preferably 0.1% or more, but when the content exceeds 3.0%, the hardenability is increased to reduce the toughness of the weld heat affected zone and the possibility of high temperature cracking in the weld heat affected zone and the weld metal. This is not desirable because there is.

The content of copper (Cu) is preferably 0.1 to 1.5%.

The Cu is an element that is effective in order to secure the base material strength and toughness due to solid solution at the base. For this purpose, Cu content should be contained more than 0.1%, but if it exceeds 1.5%, it is not preferable because it increases the hardenability by increasing the hardenability in the weld heat affected zone and promotes high temperature crack in the weld heat affected zone and weld metal. Can not do it. In particular, Cu is an element that affects the formation of acicular and polygonal ferrites effective in improving the toughness of the weld heat affected zone by depositing CuS around the Ti-based oxide with sulfur, so that the content is 0.3 to 1.5%. desirable.

In addition, in the case of complex addition of Cu and Ni, the sum thereof is preferably made less than 3.5%. The reason is that less than 3.5% of the hardenability increases, which adversely affects the weld heat affected zone toughness and weldability.

Molybdenum (Mo) is preferably 0.05 to 1.0%.

Mo is also an element that increases the hardenability and at the same time improves the strength, the content is 0.05% or more to secure the strength, but in order to suppress the hardening of the weld heat affected zone and the low temperature cracks of the weld, the upper limit is 1.0% as Cr do.

Chromium (Cr) is preferably 0.05 to 1.0%.

The Cr increases the hardenability and also improves the strength. If the content is less than 0.05%, the target strength cannot be obtained and if the content exceeds 1.0%, the base metal and the weld heat affected zone toughness may be caused.

In addition, in the present invention, at least one of Ca and REM may be further added to suppress grain growth of the old austenite.

Calcium (Ca) is preferably 0.0005 to 0.005%, and REM is preferably 0.005 to 0.05%.

The Ca and REM form an oxide having excellent high temperature stability to inhibit the growth of the austenite grains when heated in the base material and to promote the ferrite transformation during the cooling process to improve the toughness of the weld heat affected zone. In addition, Ca has the effect of controlling the coarse MnS shape during steelmaking. To this end, it is preferable to add more than 0.0005% of calcium (Ca) and more than 0.005% of REM. However, when Ca is 0.005% and more than 0.05% of REM, large inclusions and clusters are generated to impair the cleanliness of the steel. As REM, 1 or more types, such as Ce, La, Y, and Hf, may be used and the said effect can be acquired.

Hereinafter, the microstructure and the precipitate in the present invention will be described in detail.

In the present invention, after the hot rolling, the steel microstructure is made of ferrite + bainite, and the size of the ferrite grains is preferably 20 µm or less. The reason is that when the grain size of the ferrite is larger than 20 μm, the average austenite grain size of the weld heat affected zone during the high heat input welding becomes 80 μm or more, which is detrimental to the weld heat affected zone toughness.

In addition, the higher the percentage of ferrite in the composite structure of ferrite + bainite, the higher the toughness and elongation of the base metal, and the ferrite is 50% or more and most preferably 70% or more.

In addition, the old austenite grains of the weld heat affected zone are greatly influenced by the size and number of precipitates distributed in the base material and the distribution when the austenite grain size of the base material is constant. In addition, since 30-40% of the nitrides distributed in the base material are welded to the base material at the time of high heat input welding (above the heating temperature of 1400 ℃ or more), the effect of inhibiting the growth of the austenite grain growth in the welding heat affected zone is reduced. There is a need for a uniform distribution of further nitrides taking into account the re-used nitrides.

In order to suppress the growth of the austenite in the weld heat affected zone, fine TiN, precipitates and Ti-Nb and Ti-V complex precipitates are uniformly distributed to suppress the Ostwald ripening phenomenon in which some precipitates are coarsened. It is important. For this purpose, the distribution of composite precipitates should be uniform with the interval of TiN, Ti-Nb and Ti-V composite precipitates of 0.5 μm or less.

In addition, it is preferable to limit the particle size and the critical number of TiN, Ti-Nb and Ti-V composite precipitates to 0.01 to 0.1 µm and 1.0 × 10 ⁷ or more per mm ² . The reason for this is that when the heat input welding is less than 0.01 μm, most of them are easily re-used in the base metal, and thus the effect of inhibiting the growth of the old austenite grains is insufficient. When the thickness exceeds 0.1 μm, pinning of the old austenite grains is performed. This is because the grain growth inhibition effect is reduced and the same behavior as coarse nonmetallic inclusions has a detrimental effect on the mechanical properties. In addition, when the number of precipitates is less than 1.0 × 10 ⁷ per 1 mm ² , it may be difficult to control the size of the old austenite grain size of the high heat input welding heat affected zone below a threshold of 80 μm.

 Hereinafter, the present invention will be described in detail by way of examples.

EXAMPLE

Steel grades having the composition as shown in Tables 1 and 2 were prepared as slabs by a continuous casting method by dissolving them in a sample, and the composition ratio between alloying elements for each steel type to show the effect of the present invention is shown in Table 2 below. .

Test specimens for evaluating the mechanical properties of the base metal as described above were rolled and collected at the center of the plate thickness of the rolled material. The tensile test specimens were taken in the rolling direction and the Charpy impact specimens were taken in the direction perpendicular to the rolling direction.

Tensile test piece was used KS standard (KS B 0801) No. 4 test piece and the tensile test was tested at the cross head speed (5mm / mim). The impact test piece was manufactured according to KS (KS B 0809) No. 3 test piece, and the notch direction was processed on the side of the rolling direction (L-T) in the case of the base material and in the welding line direction on the welding material. In addition, the high temperature toughness evaluation of the base material measured the reduction rate in 1000 degreeC.

In addition, in order to investigate the austenite grain size according to the maximum heating temperature of the welding heat affected zone, it is heated to 140 ℃ / sec condition for 1 second after the maximum heating temperature (1200 ~ 1400 ℃) using the simulation welding simulator (simulator). After holding, it was quenched with He gas. The quenched specimens were ground and corroded to determine the austenite grain size at the highest heating temperature condition by KS (KS D 0205).

The size, number, and spacing of precipitates and oxides, which have a significant effect on the microstructure analysis and the toughness of the weld affected zone after cooling, were measured by the point counting method using an image analyzer and an electron microscope. At this time, the test surface was evaluated based on 100 mm ² .

Impact toughness evaluation of the welding heat affected zone is 800-500 ℃ cooling after heating the welding conditions corresponding to about 80 kJ / cm, 150 kJ / cm, 250 kJ / cm, that is, the maximum heating temperature to 1400 ℃ After the welding heat cycles of 60 seconds, 120 seconds, and 180 seconds were given, the surface of the test piece was polished, processed into an impact test piece, and evaluated through a Charpy impact test at -40 ° C. In addition, in order to evaluate the impact toughness of the actual high heat input welding joint, the electrogas welding was performed with a welding heat input equivalent to about 600 kJ / cm, and evaluated by Charpy impact test at -20 ° C in the fusion line.

Chemical composition (% by weight) C Si Mn Ti Al N B W Nb V P S O (ppm) (ppm) Inventive Steel 1 0.12 0.13 1.54 0.014 0.04 115 7 0.005 0.01 0.01 0.006 0.005 20 Inventive Steel 2 0.1 0.12 1.48 0.02 0.02 80 5 0.001 0.05 0.01 0.006 0.005 11 Invention Steel 3 0.1 0.14 1.5 0.02 0.025 100 0 0.004 0.09 0.03 0.007 0.005 12 Inventive Steel 4 0.13 0.14 1.48 0.015 0.04 115 8 0.15 0.02 0.01 0.007 0.005 10 Inventive Steel 5 0.13 0.21 1.5 0.02 0.025 90 4 0.002 0.02 0.01 0.007 0.005 13 Inventive Steel 6 0.07 0.16 1.45 0.025 0.045 100 6 0.05 0.02 0.01 0.008 0.006 12 Comparative Steel 1 0.08 0.15 1.52 0.05 0.09 300 15 0.002 0.05 0.04 0.006 0.004 12 Comparative Steel 2 0.11 0.15 1.52 0.018 0.06 120 10 0.001 0.01 0.015 0.007 0.005 14 Comparative Steel 3 0.09 0.12 1.48 0.019 0.048 130 10 0.01 0.01 0.01 0.006 0.003 12 Comparative Steel 4 0.07 0.12 1.5 0.05 0.07 275 10 0.002 0.01 0.02 0.006 0.005 24 Comparative Steel 5 0.14 0.1 1.48 0.015 0.06 112 3 0.003 0.02 0.01 0.006 0.005 18 Conventional Steel 1 0.05 0.13 1.31 0.009 0.0014 22 1.6 - - - 0.002 0.006 22 Conventional Steel 2 0.05 0.11 1.34 0.012 0.0036 48 0.5 - - - 0.002 0.003 32 Conventional Steel 3 0.13 0.24 1.44 0.01 0.0044 127 1.2 - 0.05 - 0.012 0.003 138 Conventional Steel 4 0.06 0.18 1.35 0.013 0.0027 32 8 - - 0.028 0.008 0.002 25 Conventional Steel 5 0.06 0.18 0.88 0.013 0.0021 20 5 - 0.015 0.037 0.006 0.002 27 Conventional Steel 6 0.13 0.27 0.98 0.009 0.001 28 11 - 0.001 0.045 0.005 0.001 25 Conventional Steel 7 0.13 0.24 1.44 0.008 0.02 79 8 - 0.036 - 0.004 0.002 - Conventional Steel 8 0.07 0.14 1.52 0.007 0.002 57 4 - 0.013 - 0.004 0.002 - Conventional Steel 9 0.06 0.25 1.31 0.007 0.019 91 10 - 0.025 0.035 0.008 0.002 - Conventional Steel 10 0.09 0.26 0.86 0.008 0.046 142 15 - 0.021 0.021 0.009 0.003 - Conventional Steel 11 0.14 0.44 1.35 0.049 0.03 89 7 - - 0.069 0.012 0.012 - Conventional steels (1, 2, 3) are inventive steels (5, 32, 55) of Japanese Patent Application Laid-open No. Hei 9-194990. Conventional steels 4, 5 and 6 are inventive steels 14, 24 and 28 of JP-A-10-298908. Conventional steels 7, 8, 9 and 10 are inventive steels 48, 58, 60 and 61 of JP-A-8-60292. Conventional steel 11 is invention steel F of JP-A-11-140582

Chemical composition (% by weight) Ni Cu Mo Cr Ca REM Ti / N N / B Nb / N V / N Al * N Inventive Steel 1 - - - - - - 1.2 16.4 0.9 0.9 4.6 × 10 ^-4 Inventive Steel 2 - - - - - - 2.5 16 6.3 1.3 1.6 × 10 ^-4 Invention Steel 3 - - 0.1 - - - 2 10 9 3 2.5 × 10 ^-4 Inventive Steel 4 - 0.1 - - - - 1.3 14.4 1.7 0.9 4.6 × 10 ^-4 Inventive Steel 5 - - - 0.1 0.001 - 2.2 22.5 2.2 1.1 2.25 × 10 ^-4 Inventive Steel 6 0.3 - - - - 0.01 2.5 16.7 2 One 4.5 × 10 ^-4 Comparative Steel 1 - 0.1 - 0.1 - - 1.7 20 1.7 1.3 27 × 10 ^-4 Comparative Steel 2 - - - - - - 1.5 12 0.8 1.3 7.2 × 10 ^-4 Comparative Steel 3 0.2 - - - - - 1.5 13 - 0.8 6.24 × 10 ^-4 Comparative Steel 4 0.2 - - - - - 1.8 27.5 0.4 0.7 19.25 × 10 ^-4 Comparative Steel 5 - 0.1 - - - - 1.3 37.3 1.8 0.9 6.72 × 10 ^-4 Conventional Steel 1 - - - - - - 4.1 13.8 - - 0.308 × 10 ^-5 Conventional Steel 2 - - - - - - 2.5 96 - - 1.728 × 10 ^-5 Conventional Steel 3 - 0.3 0.05 - - - 0.8 105.8 3.9 - 5.588 × 10 ^-5 Conventional Steel 4 - - - 0.14 - - 4.1 4 - 8.8 0.864 × 10 ^-5 Conventional Steel 5 0.58 0.75 0.015 0.24 - - 6.5 4 7.5 18.5 0.42 × 10 ^-5 Conventional Steel 6 1.15 0.35 0.001 0.53 - - 3.2 2.5 0.4 16.1 0.28 × 10 ^-5 Conventional Steel 7 - 0.3 0.036 - - - One 9.9 4.6 - 1.58 × 10 ^-4 Conventional Steel 8 0.35 0.32 0.013 - - - 1.2 14.3 2.3 - 1.14 × 10 ^-5 Conventional Steel 9 - - 0.025 0.21 - - 0.8 9.1 2.7 3.8 1.729 × 10 ^-4 Conventional Steel 10 1.09 - 0.36 0.51 - - 0.6 9.5 1.5 1.5 6.532 × 10 ^-4 Conventional Steel 11 - - - - - - 5.5 12.7 - 7.8 2.67 × 10 ^-4

division TiN + (Ti-Nb) N + (Ti-V) N precipitates Base material texture characteristics Base material mechanical properties High temperature ductility The number (pieces / mm ²⁾ Average size (㎛) Thickness (㎛) FGS (μm) Ferrite Percentage (%) Thickness (mm) Yield strength (MPa) Tensile strength (MPa) Elongation (%) -40 ℃ impact toughness (J) 1000 ℃ reduction rate (%) Inventive Steel 1 3.2 × 10 ⁸ 0.025 0.35 6 83 30 410 588 35 359 54 Inventive Steel 2 4.3 × 10 ⁸ 0.014 0.35 5 82 30 425 583 36 352 65 Invention Steel 3 5.2 × 10 ⁸ 0.021 0.35 7 82 40 420 587 35 358 57 Inventive Steel 4 3.7 × 10 ⁸ 0.029 0.29 9 84 40 412 597 36 362 58 Inventive Steel 5 3.2 × 10 ⁸ 0.024 0.34 6 81 40 426 568 36 362 56 Inventive Steel 6 3.2 × 10 ⁸ 0.025 0.35 6 85 40 410 558 35 350 60 Comparative Steel 1 5.6 × 10 ⁸ 0.028 0.36 5 81 30 435 587 36 348 12 Comparative Steel 2 3.2 × 10 ⁸ 0.025 0.25 7 83 40 412 583 35 347 21 Comparative Steel 3 3.2 × 10 ⁸ 0.023 0.36 7 84 40 409 563 38 364 22 Comparative Steel 4 3.3 × 10 ⁸ 0.026 0.42 6 84 30 423 584 35 330 15 Comparative Steel 5 4.6 × 10 ⁸ 0.024 0.45 5 83 30 420 585 35 346 23 Conventional Steel 1 35 406 436 - Conventional Steel 2 35 405 441 - Conventional Steel 3 25 629 681 - Conventional Steel 4 Precipitates of MgO-TiN 3.03 × 10 ⁶ pcs / mm ² 40 472 609 32 Conventional Steel 5 Precipitates of MgO-TiN 4.07 × 10 ⁶ pcs / mm ² 40 494 622 32 Conventional Steel 6 Precipitates of MgO-TiN 2.80 × 10 ⁶ pcs / mm ² 50 812 912 28 Conventional Steel 7 25 629 681 - Conventional Steel 8 50 504 601 - Conventional Steel 9 60 625 648 - Conventional Steel 10 60 760 829 - Conventional Steel 11 0.2μm or less 11.1 × 10 ³ 50 401 514 18.3

As shown in Table 3, the number of precipitates (TiN and Ti-Nb and Ti-V composite precipitates) of the hot rolled material manufactured using the invention steel (1-6) satisfying the component range of the present invention is 2 × 10 ⁸ pcs / mm ² or more, while conventional steel (4-6) is 4.07 × 10 ⁶ pcs / mm ² It can be seen that the number of the invention steel is significantly increased while having a fairly uniform and fine precipitate size compared to the conventional steel below.

On the other hand, in the case of the invention steel (1-6) that satisfies all the Al * N range aimed by the present invention secured excellent high-temperature ductility as the area reduction rate when processing at 1000 ℃ 54-65%. Therefore, it can be seen that the invention steel (1-6) is greatly reduced in defects on the surface of the steel material is possible to commercialize the excellent quality.

On the contrary, in the case of the comparative steel (1-5) which does not satisfy the Al * N range aimed at by the present invention, it exhibited hot ductility for heat with an area reduction rate of 12 to 23% at 1000 ° C.

 division Austenitic grain size of welding heat affected zone (㎛) Microstructure of welding heat affected zone with 100kJ / cm heat input Reproduction Weld Heat Affected Zone -40 ℃ Impact Toughness (J) (Maximum Heating Temperature: 1400 ℃) 1200 (℃) 1300 (℃) 1400 (℃) Ferrite Percentage (%) Ferrite Average Grain Size (μm) Δt 800-500 = 60 seconds Δt 800-500 = 120 seconds Δt 800-500 = 180 seconds Impact Toughness (J) Transition Temperature (℃) Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Inventive Steel 1 23 33 55 75 9 373 -734 332 -68 294 -64 Inventive Steel 2 22 34 55 78 10 385 -76 350 -69 303 -65 Invention Steel 3 25 35 62 77 10 354 -72 328 -68 285 -65 Inventive Steel 4 26 40 56 75 10 364 -72 334 -68 296 -63 Inventive Steel 5 24 31 54 78 10 366 -73 337 -67 293 -64 Inventive Steel 6 27 30 52 75 10 363 -73 339 -68 297 -63 Comparative Steel 1 65 120 221 35 26 105 -40 30 -32 19 -25 Comparative Steel 2 25 30 52 76 9 384 -73 354 -69 304 -64 Comparative Steel 3 24 32 50 76 9 367 -74 330 -68 285 -64 Comparative Steel 4 22 34 55 78 10 385 -76 350 -69 303 -65 Comparative Steel 5 23 32 56 76 11 367 -73 330 -68 295 -64 Conventional Steel 1 -58 Conventional Steel 2 -55 Conventional Steel 3 -54 Conventional Steel 4 230 93 132 (0 ℃) Conventional Steel 5 180 87 129 (0 ℃) Conventional Steel 6 250 47 60 (0 ℃) Conventional Steel 7 -60 -61 Conventional Steel 8 -59 -48 Conventional Steel 9 -54 -42 Conventional Steel 10 -57 -45 Conventional Steel 11 219 (0 ℃)

Table 4 shows the properties of the weld heat affected zone of the inventive steel (1-6), comparative steel (1-5) and conventional steel (1-11). As shown in Table 4, the austenitic grain size of the inventive steel (1-6) at the highest heating temperature of 1400 ℃ conditions, such as the welding heat affected zone has a range of 52 ~ 62㎛, therefore, in the present invention steel It can be seen that the effect of suppressing the austenite grains in the weld heat affected zone during welding is very good.

In addition, the ferrite phase fraction of the inventive steel (1-6) in the weld heat affected zone corresponding to the weld heat input amount of 100 kJ / cm is about 75% or more, thereby ensuring excellent impact toughness of the high heat input welded heat affected zone.

As described above, according to the present invention, by limiting the amount of AlN which can cause defects on the surface of the steel product, it is possible to improve the high temperature ductility of the steel, and to significantly reduce the defect rate of the product by drastically reducing the defects on the steel surface. have. In addition, it is possible to provide a welded structural steel plate having excellent high heat input welding heat affected zone toughness.

Claims (3)

By weight C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.005-0.05%, N: 0.005-0.020%, B: 0.0003-0.01 Remaining Fe and other unavoidable impurities, including%, W: 0.001-0.2%, Nb: 0.005-0.1%, V: 0.005-0.1%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less And Ti, N, B, Nb, V, and Al are 1.2 ≦ Ti / N ≦ 4, 3 ≦ N / B ≦ 40, 0.3 ≦ Nb / N ≦ 9, 0.5 ≦ V / N ≦ 7, 8 × A thick sheet steel for high heat ductility welding, which satisfies 10 ⁻⁵ ≦ Al * N ≦ ⁵ × 10 ^−4, and is composed of a composite structure of ferrite and bainite having a microstructure of the base material of 20 μm or less. The high temperature ductility of claim 1, wherein the steel further includes at least one of Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%. This excellent plate steel for high heat input welding. The thick plate steel for high heat ductility welding of claim 1 or 2, wherein the steel further comprises at least one of Ca: 0.0005 to 0.005% and REM: 0.005 to 0.05%.