KR20210009934A - Steel plate with superior HAZ toughness for high heat input welding and method for the same - Google Patents

Abstract

Provided are structural steel having enhanced toughness of a high heat-input welding heat-affected zone and a manufacturing method thereof. The structural steel having enhanced toughness of a high heat-input welding heat-affected zone comprises, by wt%, C: 0.03 to 0.19%, Si: 0.01 to 0.5%, Mn: 0.4 to 2.0%, Ti: 0.005 to 0.05%, Al: 0.08 to 0.5%, N: 0.002 to 0.008%, O: 0.001 to 0.008%, B: 0.0003 to 0.005%, P: 0.03% or less, S: 0.03% or less (where 2.5 <= Ti/N <= 5, 20 <= Al/N <= 100, 60 <= (10Ti+Al)/N <= 125, 52 <= (C+Si+Mn/8)/N <= 120), and the remainder of Fe and other impurities, and the matrix microstructure is formed as a composite structure of ferrite and pearlite having a grain size of 20 μm or less.

Description

Steel plate with superior HAZ toughness for high heat input welding and method for the same}

The present invention relates to a structural steel used in welded structures such as architecture, bridges, shipbuilding, offshore structures, steel pipes, and line pipes. More specifically, the present invention relates to a structural steel for welding that can finely disperse TiN and AlN precipitates and suppress or refine the formation of M-A to improve the toughness of the heat-affected zone with high heat input, and a manufacturing method thereof.

Recently, in accordance with the trend of high-rise buildings and structures, the steel materials used have become large and are being replaced by thick steel materials. In order to weld such a thick material, high-efficiency welding is inevitable. As a technology for welding thickened steel, a high heat input submerged welding method and an electro-gas welding method capable of one-pass welding are widely used. In addition, in the field of shipbuilding and bridges, in order to improve welding productivity even when welding steel materials having a thickness of 25 mm or more, the high heat input welding method capable of one-pass welding as described above is largely applied.

In general, in welding, the larger the amount of heat input, the larger the amount of welding, the greater the number of welding passes, and therefore, high heat input welding is advantageous in terms of welding productivity. That is, if the amount of heat input in welding is increased, the range of use can be widened. The range of high heat input currently being used is approximately 100-300kJ/cm, but in order to weld steel with a thickness of 50mm or more thickened in one pass, it is possible to have a super heat input range of 300-600kJ/cm.

However, the heat-affected zone formed during welding, especially the heat-affected zone near the fusion boundary, is heated to a temperature close to the melting point by the amount of heat input from the steel. Since austenite grains grow and coarsen, and microstructures weak in toughness such as upper bainite and martensite are formed in the cooling process, the toughness of the heat-affected zone for welding is greatly reduced.

Therefore, in order to secure the stability of the welded structure, it is necessary to suppress the growth of austenite grains in the weld heat-affected zone and keep it fine. As a means of solving this problem, a technique for delaying the growth of crystal grains of a heat-affected zone during welding by appropriately distributing oxides or Ti-based carbonitrides, etc., which are stable at high temperatures, in a steel material is disclosed. For example, the invention described in Patent Document 1-11 and the technology described in Non-Patent Document 1 can be cited.

Among them, Patent Document 1 is a representative technology that uses TiN precipitates, and when a heat input of 100J/cm (maximum heating temperature 1400℃) is applied, a structural steel material with an impact toughness of about 200J (a base material is about 300J) at 0℃ It is disclosed. In this patent document 1, Ti/N is substantially managed as 4 to 12, and TiN precipitates of 0.05 µm or less are 5.8 × 10 ³ /mm2 to 8.1 × 10 ⁴ /mm 2, and the TiN precipitates of 0.03 to 0.2 µm are 3.9 Precipitating at a rate of ³ × 10 ³ /mm 2 to ⁴ × 10 ⁴ /mm 2 to make ferrite finer to secure the toughness of the weld.

However, according to this Patent Document 1, when a high heat input welding of 100 kJ/cm is applied, the toughness of the base material and the heat-affected zone is generally low (base material: 320J, the heat-affected zone: 220J with the highest impact toughness of 0°C), and, The difference in toughness between the base material and the heat-affected zone is about 100J, so there is a limit to securing the reliability of the steel structure by superheat input welding of thickened steel. In addition, as a method of securing the desired TiN precipitate, the slab is rapidly cooled by heating at a temperature of 1050°C or higher, and then reheated for hot rolling. do.

On the other hand, Patent Document 4 is a technology using a composite oxide composed of Al, Mn and Si by managing the ratio of Al and O to 0.3≤Al/O≤1.5 in low nitrogen steel (N≤0.005%), but about 100kJ/cm When high heat input welding is applied, the transition temperature of the heat-affected zone is -50℃, so toughness is not good. In addition, Patent Document 2 is a technique using MgO-TiN composite precipitates, but when high heat input welding of about 100 kJ/cm is applied, the impact toughness of the weld heat affected zone at 0° C. is 130 J, which is not good.

Until now, there are many known technologies that improve the toughness of the welding heat-affected zone during high heat input welding by using TiN precipitates and Ti, Al, and Mg-based oxides, but the welding heat effect during ultra-heat input welding maintained at a high temperature of 1350℃ or higher for a long time. No case has yet been published that dramatically improved the personality of wealth. In particular, there are few technologies in which the toughness of the high heat input welding heat-affected zone is similar to that of the base metal. Therefore, if the above problems can be solved, it is possible to achieve super heat input welding of thickened steel materials, thereby enhancing the efficiency of welding work and securing the reliability of the steel structure for welding by securing the toughness of the welded part of the steel structure.

Japanese Patent Laid-Open Publication No. 11-140582 Japanese Patent Laid-Open Publication No. Hei 10-298708 Japanese Patent Laid-Open Publication No. Hei 10-298706 Japanese Patent Laid-Open Publication No. Hei 9-194990 Japanese Patent Application Laid-Open No. Hei 9-324238 Japanese Patent Application Laid-Open No. Hei 8-60292 Japanese Laid-Open Patent Publication (Small) 60-245768 Japanese Patent Laid-Open Publication No. Hei 5-186848 Japanese Laid-Open Patent Publication (Small) 58-31065 Japanese Unexamined Patent Publication (Small) 61-79745 Japanese Patent Laid-Open Publication No. 64-15320.

Journal of the Japanese Welding Society Vol. 52, No. 2, page 49

An object of the present invention is to provide a steel for a welded structure having excellent toughness in a high heat input welding heat-affected zone and a manufacturing method thereof by suppressing the formation of fine TiN and AlN precipitates and M-A.

In addition, the technical problems to be achieved in the present invention are not limited to the technical problems mentioned above, and other technical problems not mentioned are clearly understood by those of ordinary skill in the technical field to which the present invention belongs from the following description. Can be.

The present invention for achieving the above object,

By weight C:0.03~0.19%, Si:0.01~0.5%, Mn:0.4~2.0%, Ti:0.005~0.05%, Al: 0.08~0.5%, N:0.002~0.008%, O:0.001~0.008 % B:0.0003~0.005%, P:0.03% or less, S:0.03% or less, 2.5≤Ti/N≤5, 20≤Al/N≤100, 60≤(10Ti+Al)/N≤125, 52≤ Welding heat consisting of a composite structure of ferrite and pearlite that satisfies (C+Si+Mn/8)/N≤120, is composed of the remaining Fe and other impurities, and has a grain size of less than 20㎛ of the parent material microstructure It relates to a welded structural steel excellent in the toughness of the affected area.

In the steel, Ni:0.1~3.0%, Cu:0.1~1.5%, Mo:0.05~1.0%, Cr:0.05~1.0%, W:0.001~0.2%, Nb: 0.005-0.1%, V:0.005~ One or more selected from the group consisting of 0.1%, and one or more selected from Ca: 0.0005 to 0.005% and REM: 0.005 to 0.05% may be additionally included.

TiN, AlN precipitates and carbides having an average size of 0.05 to 0.09 μm may be distributed in the steel material at an interval of 1.0 μm or less and 1.0×10 ⁷ pieces/mm 2 or more.

In addition, the present invention,

By weight C:0.03~0.19%, Si:0.01~0.5%, Mn:0.4~2.0%, Ti:0.005~0.05%, Al: 0.08~0.5%, N:0.002~0.008%, O:0.001~0.008 % B:0.0003~0.005%, P:0.03% or less, S:0.03% or less, 2.5≤Ti/N≤5, 20≤Al/N≤100, 60≤(10Ti+Al)/N≤125, 52≤ (C+Si+Mn/8)/N≤120, and after heating the steel slab composed of the remaining Fe and other impurities in the range of 1000-1250℃ for 60 to 180 minutes, rolling over 40% in the austenite recrystallization zone The present invention relates to a method for manufacturing structural steel having excellent toughness in a welded heat-affected zone comprising hot rolling at a ratio and then cooling at a cooling rate of 10°C/sec or less from the rolling end temperature to room temperature.

In the steel slab, Ni:0.1~3.0%, Cu:0.1~1.5%, Mo:0.05~1.0%, Cr:0.05~1.0%, W:0.001~0.2%, Nb:0.005-0.1%, V:0.005 One or more selected from the group consisting of ~0.1%, and at least one selected from among Ca: 0.0005~0.005% and REM: 0.005~0.05% may be additionally included.

In addition, the present invention,

High heat input welding is applied to the steel material (base material) manufactured by the above-described manufacturing method to create a prior austenite of 120 μm or less in the weld heat-affected area, followed by rapid cooling, resulting in a microstructure of the weld heat-affected area of 20 μm. It relates to a welded structure in which the following ferrite constitutes a phase fraction of 70% or more.

The ferrite may be at least one of polygonal ferrite and acicular ferrite structures.

The present invention having the configuration as described above, in the development of a structural steel for welding that can secure the properties of the high heat input welding heat-affected zone while having excellent base material properties, the effect of high heat input welding heat by using fine TiN and AlN precipitates. By controlling austenite grains in the part, promoting polygonal ferrite transformation within the grains, reducing solid solution nitrogen and reducing MA content, it is possible to secure excellent high heat input welding heat-affected zone toughness at the same time, and effectively provide welded structural steel with low manufacturing cost. can do.

Hereinafter, the present invention will be described.

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

As a result of investigating the grain size of the old austenite that affects the toughness of the heat-affected zone, the present inventors found that the toughness of the heat-affected zone changes based on the grain size (about 150 μm) of the critical old austenite.

Based on these studies, in the present invention,

[1] With the use of TiN precipitates and AlN complex precipitates,

[2] When high heat input welding is applied with the initial ferrite grain size of steel below the critical level, the old austenite in the heat affected zone is reduced to an average of 150㎛ or less. Also,

[3] The amount of M-A (Martensite-Austenite) can be significantly reduced by suppressing or miniaturizing the formation of carbides during the cooling process of the old austenite in the weld heat affected zone.

[4] In addition, BN and AlN precipitates are used to reduce the content of soluble nitrogen in the cooling process of the heat-affected zone.

These [1][2][3][4] will be described in more detail.

[1] TiN precipitate and AlN, BN complex precipitate management

In the case of applying high heat input welding to structural steel, the welding heat-affected area near the melting line is heated to a high temperature of about 1400℃ or higher, so that the TiN precipitate deposited in the base material is partially dissolved by the welding heat, or the Ostwald ripening phenomenon , Small precipitates are decomposed and diffused into large precipitates, and large precipitates become larger), and some precipitates are decomposed or some precipitates become coarse, and the number of TiN precipitates is remarkably reduced, leading to the growth of old austenite grains. The inhibitory effect disappears.

The present inventors focused on the fact that this phenomenon is caused by the diffusion of the solid solution Ti atoms decomposed by the heat of welding TiN precipitates distributed in the base metal, and as a result of examining the high temperature stability of the TiN precipitates according to the Ti/N ratio, that is, Ti and When the ratio of N (Ti/N) is in the range of 2.5 to 5, the amount of solid solution Ti is extremely reduced, and the high temperature stability of TiN precipitates is greatly improved, resulting in 1.0×10 ⁷ fine TiN precipitates of 0.01 to 0.1 μm size. Important results were obtained that were distributed over /mm2. This means that if the same Ti content contains an appropriate amount of nitrogen, all Ti atoms are easily combined with nitrogen atoms. In the present invention, in view of the fact that aging due to the presence of solid solution N can be promoted, the Al/N ratio, and N and Ti+Al are collectively managed so that N is compositely precipitated into TiN and AlN.

[2] Ferrite particle size management of steel (base material)

According to the research of the present invention, in order to make the average size of old austenite 150 µm in the weld heat affected zone, it is necessary to finely reduce the size of ferrite to less than about 20 µm in the base structure of ferrite + pearlite together with the management of precipitates. I confirmed that it was important. The refinement of ferrite is obtained by suppressing the growth of ferrite grains generated during the cooling process by using fine carbides (Fe ₃ C, NbC, VC, WC) together with austenite grain refinement by strong processing during hot rolling.

[3] Reduction of M-A (island martensite) content in weld heat affected zone

The findings of the present invention have been found to be effective in reducing the MA content by suppressing or minimizing the formation of carbides during the cooling process in the welding heat-affected zone during high heat input welding when the Al content in a solid solution state in the base material increases. . Therefore, by reducing the amount of M-A, it contributes to the improvement of the toughness of the high heat input welding heat-affected zone.

[4] Reduction of free nitrogen content in welding heat affected zone

It was found in the research of the present invention that during high heat input welding, part of the precipitates dispersed in the base metal is decomposed by welding heat, increasing the dissolved nitrogen content in the welding heat-affected zone, thereby reducing the toughness of the welding heat-affected zone. Likewise, the formation of TiN precipitates and AlN composite precipitates contributes to the improvement of toughness in the heat-affected zone by reducing the aging effect of the solid-solution nitrogen by reducing the content of dissolved nitrogen in the heat-affected zone.

Hereinafter, the present invention will be described in detail by dividing it into a component of a steel material and a manufacturing method thereof.

[Welded structural steel]

-The content of carbon (C) is preferably 0.03 to 0.19%.

If the content of carbon (C) is less than 0.03%, it is not preferable because it is necessary to add other alloying elements of high steel to secure strength as structural steel. In addition, when the C exceeds 0.19%, the microstructures vulnerable to toughness such as upper bainite, martensite and degenerate pearlite are transformed during cooling, reducing the low-temperature impact toughness of the structural steel, and also cracking the weld. It is not preferable due to increased sensitivity.

It is preferable to limit the content of silicon (Si) to 0.01 to 0.5%.

If the content of silicon is less than 0.01%, the deoxidation effect of molten steel is insufficient during the steelmaking process and the corrosion resistance of the steel is lowered, and if it exceeds 0.5%, the effect is saturated, and when cooling after rolling, the deoxidation effect is insufficient. It is not preferable because it promotes the transformation of martensite, lowers the low-temperature impact toughness, and affects the weld crack susceptibility.

It is preferable to limit the content of manganese (Mn) to 0.4 to 2.0%.

Mn is an effective element that improves deoxidation, weldability, hot workability, and strength in steel. It is precipitated in the form of MnS + VN around Ti-based oxides and affects the formation of needle-shaped and polygonal ferrite, which is effective for improving the toughness of the heat-affected zone. It is an element that affects. Such Mn forms a substitutional solid solution in the matrix structure to solid-solution strengthen the matrix to secure strength and toughness. For this purpose, it is recommended to add 0.4% or more. However, when the Mn content exceeds 2.0%, the structure inhomogeneity caused by Mn segregation has a detrimental effect on the toughness of the heat affected zone rather than the solid solution strengthening effect. In addition, when Mn content is added more than 2.0%, macro and micro segregation occurs according to the segregation mechanism during solidification of the steel, which promotes the formation of central segregation zones in the center during rolling, and can act as a cause of creating a low-temperature transformation structure in the center of the base material. have.

It is preferable to limit the content of aluminum (Al) to 0.08 to 0.5%.

Al is an indispensable element in the present invention because it reacts with oxygen to form a fine Al ₂ O ₃ composite oxide that is stable at high temperatures. In addition, Al not only forms fine AlN precipitates in the steel, but also plays a role in reducing MA in the heat affected zone. In order to form fine AlN precipitates and to reduce the MA content, it is necessary to increase the Al content to 0.08% or more. However, if it exceeds 0.5%, the generation of Widmanstatten ferrite and island martensite, which are weak in toughness in the cooling process of the welding heat-affected zone, may be encouraged, and rather, the toughness of the high-heat input welding heat-affected zone may be reduced.

It is preferable to limit the content of titanium (Ti) to 0.005 to 0.05%.

Ti is an indispensable element in the present invention because it combines with N to form fine TiN precipitates that are stable at high temperatures. In order to obtain such a fine TiN precipitate effect, it is preferable to add more than 0.005% of Ti, but if it exceeds 0.05%, coarse TiN precipitates are formed in the molten steel and are mixed in the cast slab and the base metal, and the weld heat affected zone during welding. It is not preferable because it cannot suppress the growth of night crystal grains.

It is preferable to limit the content of boron (B) to 0.0003 to 0.005%.

B is an element that is very effective in generating polygonal ferrite at grain boundaries as well as acicular ferrite having excellent toughness within crystal grains. B forms BN precipitates to interfere with the growth of old austenite grains, and form Fe carbides in grain boundaries and grains to promote needle-like and polygonal ferrite transformation with excellent toughness. If the B content is less than 0.0003%, such an effect cannot be expected, and if it exceeds 0.005%, the hardening property increases, which is not preferable because there is a possibility of hardening and low temperature cracking of the weld heat affected zone.

It is preferable to limit the content of nitrogen (N) to 0.002-0.008%.

N is an indispensable element for forming TiN, BN, and AlN, and during high heat input welding, it suppresses the growth of old austenite grains in the heat-affected zone to the maximum and increases the amount of precipitates such as TiN, AlN, and BN. Particularly, since the size of the TiN and AlN precipitates and the spacing of the precipitates, the distribution of the precipitates, the frequency of complex precipitation with oxides, and the high temperature stability of the precipitates themselves, the content is preferably set to 0.002% or more. However, if the nitrogen content exceeds 0.008%, the effect is saturated, and the toughness decreases due to the increase in the amount of dissolved nitrogen distributed in the weld heat affected zone, and it is mixed into the weld metal due to dilution during welding, resulting in a decrease in the toughness of the weld metal. I can.

-It is preferable to limit the content of oxygen (O) to 0.001-0.008%.

O is an essential element that reacts with Al to form an Al ₂ O ₃ composite oxide.If the O content is less than 0.001%, it is not preferable because it cannot form a fine Al ₂ O ₃ composite oxide.If it exceeds 0.008%, it is not preferable. _It is not preferable because oxides such as ₂ O ₃ composite oxide and FeO are generated, which adversely affects the properties of the heat-affected zone.

It is preferable to limit the content of phosphorus (P) and sulfur (S) to 0.030% or less, respectively.

Since P is an impurity element that promotes central segregation during rolling and high temperature cracking during welding, it is desirable to manage it as low as possible. In order to improve the toughness of the base material, the toughness of the heat-affected zone, and reduce central segregation, it is recommended to manage it below 0.03%.

When S is present in a large amount, it is preferable to manage it as low as possible because it forms a low melting point compound such as FeS. In order to reduce the base material toughness, the toughness of the heat-affected zone, and central segregation, it is recommended that the S content be less than 0.03%. In particular, in the case of S, since precipitation in the form of MnS around Ti-based oxides affects the generation of needle-shaped and polygonal ferrite, which is effective in improving the toughness of the welding heat-affected zone, the more preferable range is 0.003 when considering high-temperature cracking during welding. It is preferable to limit it to% or less.

It is preferable that the ratio of Ti/N is 2.5 to 5.

In the present invention, the Ti/N ratio is set to 2.5 or more. If the Ti/N ratio is greater than 2.5, all the solid-solution Ti atoms are stoichiometrically bonded with nitrogen atoms during the cooling process during the playing process, increasing the amount of fine TiN precipitation. In addition, stability at high temperatures increases in the case of TiN precipitates even at a high temperature such as a welding heat affected zone. On the other hand, when the Ti/N ratio is higher than 5, coarse TiN is crystallized in the molten steel during the steelmaking process, so that a uniform distribution of TiN is not obtained. In addition, the excess Ti that does not precipitate as TiN exists in a solid solution state, so welding heat It adversely affects the toughness of the affected area. And if the Ti/N ratio is less than 2.5, it may cause difficulties in controlling the quality of the slab surface during playing.

It is preferable to set the ratio of Al/N to 20-100.

In the present invention, when the Al/N ratio is less than 20, the distribution of AlN precipitates for inducing polygonal ferrite transformation is insufficient, and there is a possibility that welding cracks may occur due to an increase in the amount of solute nitrogen in the weld heat-affected zone, and the Al/N ratio exceeds 100. If so, the effect is saturated.

It is preferable that the ratio of (10Ti+Al)/N is 60 to 125.

In the present invention, when the (10Ti+Al)/N ratio is less than 60, the distribution of Al ₂ O ₃ composite oxide, Ti precipitate, and AlN precipitate for inducing polygonal ferrite transformation is insufficient, and the (10Ti+Al)/N ratio is 125. If it exceeds, it is not preferable because coarse Al ₂ O ₃ composite oxides, Ti precipitates and AlN precipitates are formed, which adversely affects the properties of the heat-affected zone.

It is preferable that the ratio of (C+Si+Mn/8)/N is 52-120.

In the present invention, when the ratio of (C+Si+Mn/8)/N is less than 52, it inhibits the growth of old austenite grains in the weld heat affected zone, generates fine polygonal ferrite at grain boundaries, and acicular and polygonal ferrites in grains When the number of fine carbides for production and control of the tissue fraction is insufficient, and (C+Si+Mn/8)/N exceeds 120, the effect is saturated.

In the present invention, to the steel formed as described above, in order to further improve mechanical properties, one or two or more selected from the group of Ni, Cu, Mo, Cr, V, Nb, and W may be additionally added.

It is preferable to limit the content of nickel (Ni) to 0.1 to 3.0%.

Ni is an effective element that improves the strength and toughness of the base metal by solid solution strengthening. In order to obtain this effect, it is preferable to contain Ni content of 0.1% or more, but if it exceeds 3.0%, it increases the hardenability to decrease the toughness of the heat-affected zone and the possibility of high-temperature cracking in the heat-affected zone and the weld metal. This is not desirable because of this.

It is preferable to limit the content of copper (Cu) to 0.1 to 1.5%.

Cu is a solid solution in the base and is an effective element to secure the base metal strength and toughness due to the solid solution strengthening effect. For this purpose, the Cu content should be 0.1% or more, but if it exceeds 1.5%, it is not preferable because it increases the hardenability in the weld heat-affected zone to reduce toughness and promotes high-temperature cracking in the weld heat-affected zone and the weld metal. . In particular, the Cu is an element that affects the formation of needle-shaped and polygonal ferrite, which is effective in improving the toughness of the welded heat-affected zone by depositing it in the form of CuS around Ti-based oxides with sulfur, so the content should be 0.3 to 1.5%. desirable.

In addition, when Cu and Ni are added in combination, the sum of them is preferably less than 3.5%. The reason for this is that when it exceeds 3.5%, the hardening property becomes large, which adversely affects the toughness and weldability of the heat affected zone.

-It is preferable that chromium (Cr) is 0.05 to 1.0%.

Cr increases the hardenability and also improves the strength. When the content is less than 0.05%, the strength cannot be obtained, and when the content exceeds 1.0%, the toughness of the base metal and the heat-affected zone is deteriorated.

-Molybdenum (Mo) is preferably 0.05 to 1.0%.

Mo is also an element that increases hardenability and improves strength at the same time. Its content is 0.05% or more to secure strength, but the upper limit is set to 1.0% like Cr in order to suppress the hardening of the heat-affected zone and low-temperature cracking of welding. .

It is preferable to limit the content of tungsten (W) to 0.001-0.2%.

Tungsten is an element that uniformly precipitates as tungsten carbide (WC) on the base material after hot rolling, effectively inhibits ferrite grain growth after ferrite transformation, and inhibits the growth of old austenite grains at the initial stage of heating of the welding heat affected zone. When the content is less than 0.001%, tungsten carbide for suppressing ferrite grain growth is less distributed during cooling after hot rolling, and when more than 0.2% is added, the effect is saturated.

It is preferable to limit the content of Nb to 0.005-0.1%.

Nb is a useful element that combines with N to form NbN precipitates to reduce solid solution nitrogen in the weld heat affected zone. In order to form fine NbN precipitates, it is preferable to add 0.005% or more of Nb content. However, if the Nb content exceeds 0.1%, it may adversely affect the mechanical properties because it increases the low-temperature transformation structure in the weld heat affected zone.

It is preferable to limit the content of V to 0.005-0.1%

V is a useful element that combines with N to form VN precipitates and promotes ferrite transformation in the austenite grains in the weld heat affected zone.In order to form fine VN precipitates, it is preferable to add a V content of 0.005% or more. 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 crack susceptibility.

In addition, in the present invention, at least one of Ca and REM may be additionally added to inhibit grain growth of guaustenite.

Ca and REM form oxides with excellent high temperature stability, suppress the growth of old austenite grains when heated in the base material, and promote ferrite transformation during cooling, thereby improving the toughness of the heat affected zone. In addition, Ca has an effect of controlling the coarse MnS shape during steelmaking. To this end, it is recommended to add more than 0.0005% of calcium (Ca) and more than 0.005% of REM, but if the Ca content exceeds 0.005% and REM exceeds 0.05%, large inclusions and clusters are generated, thereby damaging the cleanliness of the steel. As REM, one or two or more of Ce, La, Y and Hf may be used, and any of the above effects can be obtained.

On the other hand, in the present invention, the microstructure of the steel material after hot rolling is ferrite + pearlite, and the size of the ferrite grains is preferably 20 μm or less. The reason is that when the grain size of ferrite is larger than 20 μm, the average grain size of the old austenite in the weld heat-affected area becomes 150 μm or more during high heat input welding, which is harmful to the toughness of the weld heat-affected area.

In the composite structure of ferrite + pearlite, as the phase fraction of ferrite increases, the toughness and elongation of the base material increase. Ferrite is 50% or more, most preferably 70% or more.

In addition, in the present invention, when the austenite grain size of the base metal is constant, the old austenite grains of the welding heat-affected zone are greatly affected by the size, number, and distribution of precipitates distributed in the base metal. In addition, in the case of the nitride distributed in the base material when welding with high heat input or higher (heating temperature of 1400°C or higher), 30 to 40% of the nitride is re-used as the base material, reducing the effect of inhibiting the growth of old austenite grains in the heat-affected zone. A uniform distribution of more nitrides is required considering the nitrides to be re-used. In order to suppress the growth of old austenite in the weld heat affected zone, it is important to uniformly distribute fine TiN precipitates and AlN and BN complex precipitates to suppress the Ostwald ripening phenomenon in which some precipitates become coarse. For this purpose, the distribution of the composite precipitates should be uniform so that the spacing of the TiN, BN and AlN composite precipitates is 1.0 μm or less.

In addition, it is preferable to limit the particle size and critical number of TiN, BN, AlN composite precipitates and fine carbides to 0.05 to 0.09 μm and 1.0×10 ⁷ or more per 1 mm ² . The reason is that if less than 0.05㎛, it is easily re-used in the base material during high heat input welding, and the effect of suppressing the growth of old austenite grains is insufficient. If it exceeds 0.09㎛, pinning for the old austenite grains This is because the effect of growth inhibition) decreases and it acts like a coarse non-metallic inclusion, which 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 is difficult to control the old austenite grain size of the high heat input welding heat-affected zone to less than 150 μm, which is a critical value.

[Method of manufacturing welded structural steel]

Hereinafter, a method of manufacturing a welded structural steel according to an embodiment of the present invention will be described.

In general, the steel refining process consists of primary refining in a converter and then out-of-furnace refining in which the molten steel in the converter is reinforced with a ladle and secondary refining.For thick materials like welded structural steels, degassing (RH process) ). Usually, deoxidation takes place between the primary and secondary refining.

Next, it is preferable to add Ti to make the content of Ti 0.005 to 0.05%, which is an advantageous range for the proper distribution of TiN precipitates for showing the effect of the present invention.

In the present invention, the molten steel refining as described above is continuously cast to form a steel slab. In view of productivity improvement, continuous casting is desirable to cast at low speed in consideration of the high probability of occurrence of cracks on the surface of cast iron in high nitrogen steel, and to provide mild cooling conditions in a secondary cooling zone. In the secondary cooling zone, the cooling conditions are an important factor affecting the micronization and uniform distribution of TiN precipitates.

According to the research of the present inventors, the continuous casting speed is less than 1.5 m/min, which is a lower speed than about 12.0 m/min, which is a typical casting speed, more preferably about 0.9 to 1.5 m/min. The reason is that when the casting speed is less than 0.9m/min, it is advantageous for cracks on the playing surface, but productivity is lowered. If it is faster than 2.0m/min, the possibility of cracking on the playing surface is high.

In addition, in the secondary cooling zone, it is recommended that the non-water amount is as weakly cooled as possible, that is, 0.3 to 0.35ℓ/kg. When the specific amount is less than 0.3ℓ/kg, it is difficult to control the appropriate size and number of TiN and AlN precipitates to show the effect of the present invention due to coarsening of TiN precipitates. In addition, when the specific water quantity exceeds 0.35ℓ/kg, the precipitation frequency of TiN and AlN precipitates is small, so it is difficult to control the number and size of TiN precipitates to show the effect of the present invention.

In the present invention, the slab is heated at 1000 to 1250°C for 120 to 360 minutes. Below 1000℃, the diffusion rate of solute atoms is small, so the number of TiN and AlN precipitates is small. If it exceeds 1250℃, TiN and AlN precipitates become coarse or partially decompose, and the number of precipitates decreases. It is not desirable. On the other hand, if the heating time is less than 120 minutes, the diffusion effect of solute atoms is insufficient, so that the effect of reducing slab segregation is small, and there is insufficient time for the solute atoms to diffuse to form precipitates, which is not preferable. In addition, when the heating time exceeds 360 minutes, the austenite grain size becomes coarse, which is not preferable in terms of work productivity.

After heating as described above, it is preferable to perform hot rolling at a rolling ratio of 40% or more in the austenite recrystallization temperature range. The austenite recrystallization zone temperature is affected by the emphasization and the previous reduction, and the austenite recrystallization zone temperature is in the range of about 800 to 1000°C in consideration of the usual reduction in the emphasis of the present invention.

After rolling as described above, air-cooling from the final rolling end temperature of 810 to 850°C to room temperature or accelerated cooling to room temperature at a rate of 10°C/sec or less to secure fine ferrite. Cooling at a cooling rate exceeding 10°C/sec is not preferable because it transforms into a martensite structure instead of fine ferrite, which adversely affects the physical properties of the steel material.

In the present invention, the casting of the steel may be performed by continuous casting or mold casting to produce a slab. At this time, if the cooling rate is fast, it is advantageous to finely disperse the precipitate, so continuous casting with a fast cooling rate is preferable. Also, for the same reason, a thin slab is advantageous. In the hot rolling process of the slab, hot charge rolling and direct rolling may be applied depending on the user's use, and various techniques such as known control rolling and controlled cooling may be applied. In addition, heat treatment may be applied to improve the mechanical properties of the hot-rolled sheet material manufactured according to the present invention. However, even if such known techniques are applied to the present invention, it is natural to interpret this as a simple change of the present invention and substantially within the scope of the technical idea of the present invention.

[Welding structure]

In addition, in the present invention, high heat input welding is applied to the steel material (base material) manufactured by the above-described manufacturing method to generate a prior austenite of 120 μm or less in the weld heat-affected zone, followed by rapid cooling to fine It is possible to provide a welded structure in which ferrite having a structure of 20 μm or less constitutes a phase fraction of 70% or more.

Here, the ferrite may be at least one of polygonal ferrite and acicular ferrite structures.

Hereinafter, the present invention will be described in detail through examples.

(Example)

River
No. Chemical composition (% by weight) C Si Mn P S Al Ti B N Nb Cu Ni Cr Mo W V Ca REM O One 0.10 0.13 1.54 0.006 0.005 0.41 0.014 7 45 - - - - - - - 0.001 - 15 2 0.09 0.12 1.50 0.006 0.005 0.32 0.02 9 55 - - 0.2 - - - - - - 24 3 0.14 0.10 1.48 0.006 0.005 0.14 0.015 3 42 - 0.1 - - - 0.005 0.01 - - 38 4 0.08 0.15 1.52 0.006 0.004 0.52 0.04 12 80 - 0.1 - 0.1 - 0.001 - - 0.001 52 5 0.10 0.10 1.50 0.007 0.005 0.08 0.02 11 40 0.01 - - - 0.1 - - - - 18 6 0.13 0.14 1.48 0.007 0.005 0.15 0.015 8 45 - 0.1 - - - - 0.01 - - 60 7 0.10 0.15 1.52 0.007 0.005 0.25 0.018 10 39 0.01 - - - - - - - - 44 8 0.15 0.16 1.45 0.008 0.006 0.42 0.022 6 71 - - 0.3 - - - - - 0.01 72 9 0.18 0.13 1.48 0.006 0.003 0.15 0.019 3 43 - - 0.2 - - - 0.01 - - 12 10 0.05 0.13 1.31 0.002 0.006 0.0014 0.009 1.6 22 - - - - - - - - - 22 11 0.05 0.11 1.34 0.002 0.003 0.0036 0.012 0.5 48 - - - - - - - - - 32 12 0.13 0.24 1.44 0.012 0.003 0.0044 0.010 1.2 127 0.05 0.3 - - - - - - - 138 13 0.06 0.18 1.35 0.008 0.002 0.0027 0.013 8 32 - - - 0.14 0.15 - 0.028 - - 25 14 0.06 0.18 0.88 0.006 0.002 0.0021 0.013 5 20 0.015 0.75 0.58 0.24 0.14 - 0.037 - - 27 15 0.13 0.27 0.98 0.005 0.001 0.001 0.009 11 28 0.001 0.35 1.15 0.53 0.49 - 0.045 - - 25 16 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 0.036 0.3 - - - - - - - - 17 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 0.013 0.32 0.35 - - - - - - - 18 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 0.025 - - 0.21 0.19 - 0.035 - - - 19 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 0.021 - 1.09 0.51 0.36 - 0.021 - - - 20 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 - - - - - - 0.069 - - - -Kang (10-12) is the invention steel (5, 32, 55) of Japanese Laid-Open Patent Publication Hei 9-194990
-Kang (13-15) is the invention steel (14, 24, 28) of Japanese Patent Laid-Open Publication No. Hei 10-298908
-Kang (16-19) is the invention steel (48, 58, 60, 61) of Japanese Laid-Open Patent Publication Hei 8-60292
-Steel 20 is invention steel F of Japanese Laid-Open Patent Publication No. Hei 11-140582

* In Table 1, the units of B, N and O are ppm

Steel grade NO. Composition ratio of alloy elements Ti/N Al/N (10Ti+Al)/N (C+Si+Mn/8)/N One Invention Lesson 1 3.1 91 122 94 2 Invention Lesson 2 3.6 58 95 72 3 Invention Lesson 3 3.6 33 69 101 4 Invention Lesson 4 2.6 32 58 52 5 Invention Lesson 5 5.0 65 115 53 6 Invention Lesson 6 5.0 20 70 97 7 Invention Lesson 7 3.3 33 67 101 8 Invention Lecture 8 4.6 64 110 113 9 Invention Lecture 9 1.5 51 66 78 10 Conventional material 10 4.1 0.6 42 155 11 Conventional material 11 2.5 0.8 26 69 12 Conventional Steel 12 0.8 0.4 8 43 13 Conventional Steel 13 4.1 0.8 42 128 14 Conventional Steel 14 6.5 1.1 66 175 15 Conventional Steel 15 3.2 0.4 36 186 16 Conventional Steel 16 1.0 2.5 13 70 17 Conventional Steel 17 1.2 0.4 13 70 18 Conventional Steel 18 0.8 2.1 10 52 19 Conventional Steel 19 0.6 3.2 9 32 20 Conventional Steel 20 5.5 3.4 58 84

The steel types having the component composition as shown in Table 1 were used as specimens and melted in a converter to form slabs by continuous casting, and the composition ratios between the alloying elements of each steel type are shown in Table 2 above. Table 3 shows the solidification rate of slabs by steel type, slab heating temperature, heating time, rolling start and end temperature, rolling reduction, and cooling rates of rolled materials manufactured with a thickness of 25 to 50 mm in the rolling process. At this time, the cumulative reduction ratio during rolling of all steel types was set to 70% or more.

Steel grade used division Heating temperature
(℃) Heating time
(min) Rolling start temperature
(℃) Rolling end
Temperature
(℃) Reduction/cumulative reduction in recrystallization zone
(%) Cooling
speed
(℃/sec) Cooling end temperature
(℃) Invention Lesson 1 Invention Example 1 1160 280 1100 810 65/80 4 Room temperature Inventive Example 2 1200 220 1160 820 65/80 5 Room temperature Invention Example 3 1230 330 1190 830 65/80 7 Room temperature Comparative Example 1 900 45 820 740 65/80 0.5 Room temperature Comparative Example 2 1350 700 1300 1080 65/80 18 Room temperature Invention Lesson 2 Invention Example 4 1160 330 1050 810 65/80 3 Room temperature Invention Lesson 3 Invention Example 5 1120 330 1060 840 65/80 5 Room temperature Invention Lesson 4 Invention Example 6 1180 350 1060 820 60/80 3 Room temperature Invention Lesson 5 Invention Example 7 1190 240 1020 810 60/80 5 Room temperature Invention Lesson 6 Invention Example 8 1120 230 1060 860 60/80 8 Room temperature Invention Lesson 7 Invention Example 9 1150 230 1020 850 60/80 8 Room temperature Invention Lecture 8 Inventive Example 10 1140 210 1010 820 60/75 6 Room temperature Invention Lecture 9 Invention Example 11 1130 230 1050 840 60/80 4 Room temperature -Conventional steel (10-12) is manufactured under the manufacturing conditions described in Japanese Laid-Open Patent Publication Hei 9-194990
-Conventional steel (13-15) is manufactured under the manufacturing conditions described in Japanese Laid-Open Patent Publication Hei 10-298908
-Conventional steel (16-19) is manufactured under the manufacturing conditions described in Japanese Laid-Open Patent Publication No. Hei 8-60292.
-Conventional steel 20 is manufactured under the manufacturing conditions described in Japanese Laid-Open Patent Publication Hei 11-140582

Test specimens for evaluating the mechanical properties of the steel from the hot-rolled plates as described above were taken from the center of the thickness of the rolled material, and the tensile test was measured in the rolling direction, and the Charpy impact test was measured in the direction perpendicular to the rolling direction. I did.

The tensile test piece was used as a test piece of KS standard (KS B 0801) No. 4, and the tensile test was performed at a cross head speed of 5 mm/mim. The impact test piece was manufactured according to the test piece of KS (KS B 0809) No. 3, and at this time, the notch direction was processed from the side (L-T) of the rolling direction in the case of steel, and the welding material was processed in the direction of the welding line. In addition, in order to investigate the size of austenite grains according to the maximum heating temperature of the welding heat-affected zone, heat up to the maximum heating temperature (1200~1400℃) at 140℃/sec using a simulated reproducible welding tester, and then for 1 second. After holding, it was quenched using He gas. The quenched test piece was polished and corroded to measure the old austenite grain size in the highest heating temperature condition according to the KS standard (KS D 0205), and the results are shown in Table 4 below.

After cooling, the size, number, and spacing of precipitates and carbides, which have an important influence on the analysis of microstructure and toughness of the weld-affected zone, were measured by a point counting method using an image analyzer and an electron microscope. Is also shown in Table 4. At this time, the test surface was evaluated based on 100 mm ² .

The evaluation of the impact toughness of the heat-affected zone is the welding conditions equivalent to about 250kJ/cm, 350kJ/cm, and 500kJ/cm equivalent to the actual welding heat input, that is, the maximum heating temperature is heated to 1400℃ and then cooled to 800~500℃. After giving a welding heat cycle of 120 seconds, 180 seconds, and 250 seconds, respectively, the surface of the test piece was polished, processed into an impact test piece, and evaluated through a Charpy impact test at -40°C, and the results are shown in Table 5.

division TiN+AlN
Precipitate and carbide Base material structure characteristics Base metal mechanical properties Total number
(Pcs/mm ² ) Average size
(㎛) interval
(㎛) FGS
(㎛) Ferrite phase fraction (%) thickness
(mm) Yield strength
(MPa) The tensile strength
(MPa) Elongation
(%) -40℃
Impact toughness
(J) Invention Example 1 1.2X10 ⁷ 0.07 0.85 8 82 25 406 554 34 335 Inventive Example 2 1.1X10 ⁷ 0.05 0.96 7 81 25 405 562 33 316 Invention Example 3 1.4X10 ⁷ 0.06 0.94 9 83 25 412 561 34 323 Comparative Example 1 1.3X10 ⁵ 0.2 2.4 27 89 25 340 480 26 216 Comparative Example 2 1.4X10 ⁵ 0.5 3.5 33 31 25 749 881 12 85 Invention Example 4 1.5X10 ⁷ 0.07 0.95 7 83 30 411 568 34 346 Invention Example 5 1.6X10 ⁷ 0.08 0.92 8 82 30 406 552 35 336 Invention Example 6 1.6X10 ⁷ 0.07 0.80 7 80 30 418 555 35 345 Invention Example 7 1.3X10 ⁷ 0.06 0.85 6 82 30 415 563 34 332 Invention Example 8 1.6X10 ⁷ 0.06 0.86 7 83 30 415 547 38 328 Invention Example 9 1.2X10 ⁷ 0.07 0.95 8 85 40 423 567 36 338 Inventive Example 10 1.2X10 ⁷ 0.07 0.85 8 81 40 414 563 36 337 Invention Example 11 1.2X10 ⁷ 0.09 0.98 7 80 40 416 558 36 322 Conventional Steel 1 35 406 436 - Conventional Steel 2 35 405 441 - Conventional Steel 3 25 629 681 - Conventional Steel 4 MgO-TiN precipitate (3.03X10 ⁶ ) 40 472 609 32 Conventional Steel 5 MgO-TiN precipitate (4.07X10 ⁶ ) 40 494 622 32 Conventional Steel 6 MgO-TiN precipitate (2.80X10 ⁶ ) 50 812 912 28 Conventional Steel 7 25 629 681 - Conventional Steel 8 50 504 601 - Conventional Steel 9 60 526 648 - Conventional Steel 10 60 760 829 - Conventional Steel 11 50 401 514 18.3

As shown in Table 4, the number of TiN precipitates and carbides in the hot-rolled material manufactured by the present invention has a range of 1×10 ⁷ pieces/mm ² or more, whereas in the case of conventional steel, 4.07×10 ⁶ pieces/ Since it shows a range of mm ² or less, it can be seen that the invention examples have a fairly uniform and fine precipitate size compared to the conventional steels, and the number of them is also significantly increased. On the other hand, in the base material structure of the inventive steel, in the case of the inventive examples, it can be seen that the ferrite grain size (FGS) is in the range of about 6 to 9 μm, which is very fine compared to the conventional steel. It can be seen that it is composed of a high ferrite fraction.

division Welding heat affected zone austenite
Grain size (㎛) 100kJ/cm heat input
Welding heat affected zone
Microstructure Reproduced welding heat affected zone
-40℃ impact toughness (J)
(Maximum heating temperature: 1400℃) 1200
(℃) 1300
(℃) 1400
(℃) ferrite
Phase fraction
(%) Ferrite average
Grain
Size (㎛) MA content
(%) Δt _800-500 =120 seconds Δt _800-500 =180 seconds Δt _800-500 =250 seconds Shock
tenacity
(J) Transition
Temperature
(℃) Shock
tenacity
(J) Transition
Temperature
(℃) Shock
tenacity
(J) Transition
Temperature
(℃) Invention Example 1 23 63 95 73 11 2 372 -734 332 -60 290 -54 Inventive Example 2 22 66 98 74 10 3 373 -76 350 -60 301 -50 Invention Example 3 23 65 110 76 11 3 357 -73 330 -62 293 -51 Comparative Example 1 54 78 185 34 24 15 115 -43 43 -30 30 -20 Comparative Example 2 65 120 221 35 26 21 102 -40 30 -22 19 -9 Invention Example 4 25 52 89 75 14 2 334 -72 318 -60 276 -52 Invention Example 5 26 49 96 76 13 3 336 -72 32 -63 256 -53 Invention Example 6 24 42 104 75 11 2 345 -73 319 -63 290 -52 Invention Example 7 27 47 113 73 12 3 353 -73 321 -63 248 -50 Invention Example 8 24 48 112 75 10 3 352 -74 312 -62 252 -50 Inventive Example 9 22 45 97 72 10 3 365 -73 324 -64 223 -52 Inventive Example 10 27 42 97 74 11 3 344 -73 321 -62 237 -53 Invention Example 11 23 62 90 72 12 2 353 -72 315 -64 232 -54 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 5 shows the properties of the welding heat-affected zone of steels of the present invention examples, comparative examples and conventional steels. Looking at the austenite grain size under the condition of the maximum heating temperature of 1400°C as in the welding heat affected zone, it can be seen that the austenite grain size is in the range of 80 to 120 μm in the example of the present invention, whereas the comparative example has a very coarse range of about 180 μm or more. . Therefore, it can be seen that the present invention steel has a very excellent effect of suppressing austenite grains in the heat-affected zone during welding. On the other hand, it can be seen that the MA content is very small in the case of the present invention in the welding heat affected zone equivalent to the welding heat input of 100 kJ/cm, and the ferrite phase fraction is composed of about 70% or more, and is superior to the conventional steel for high heat input welding. It shows the impact toughness of the heat affected zone.

As described above, in the detailed description of the present invention, preferred embodiments of the present invention have been described, but those of ordinary skill in the art to which the present invention pertains, various modifications within the limit not departing from the scope of the present invention. Of course this is possible. Therefore, the scope of the present invention is limited to the described embodiments and should not be determined, and should not be determined by the claims to be described later, as well as those equivalent thereto.

Claims (7)

By weight C:0.03~0.19%, Si:0.01~0.5%, Mn:0.4~2.0%, Ti:0.005~0.05%, Al: 0.08~0.5%, N:0.002~0.008%, O:0.001~0.008 % B:0.0003~0.005%, P:0.03% or less, S:0.03% or less, 2.5≤Ti/N≤5, 20≤Al/N≤100, 60≤(10Ti+Al)/N≤125, 52≤ Welding heat consisting of a composite structure of ferrite and pearlite that satisfies (C+Si+Mn/8)/N≤120, is composed of the remaining Fe and other impurities, and has a grain size of less than 20㎛ of the parent material microstructure Welded structural steel with excellent toughness in the affected area.
According to claim 1, wherein the steel material, Ni: 0.1 ~ 3.0%, Cu: 0.1 ~ 1.5%, Mo: 0.05 ~ 1.0%, Cr: 0.05 ~ 1.0%, W: 0.001 ~ 0.2%, Nb: 0.005-0.1 Welding heat, characterized in that one or more selected from the group consisting of %, V: 0.005 to 0.1%, and at least one selected from Ca: 0.005 to 0.005% and REM: 0.005 to 0.05% are additionally included Welded structural steel with excellent toughness in the affected area.
According to claim 1, wherein the welding heat affected zone toughness, characterized in that the steel has been TiN, AlN precipitates and carbides having an average size of 0.05 ~ 0.09㎛ distribution above 1.0 × 10 ⁷ gae / ㎟ at intervals of less than 1.0㎛ Excellent welded structural steel.
By weight C:0.03~0.19%, Si:0.01~0.5%, Mn:0.4~2.0%, Ti:0.005~0.05%, Al: 0.08~0.5%, N:0.002~0.008%, O:0.001~0.008 % B:0.0003~0.005%, P:0.03% or less, S:0.03% or less, 2.5≤Ti/N≤5, 20≤Al/N≤100, 60≤(10Ti+Al)/N≤125, 52≤ (C+Si+Mn/8)/N≤120, and after heating the steel slab composed of the remaining Fe and other impurities in the range of 1000-1250℃ for 60 to 180 minutes, rolling over 40% in the austenite recrystallization zone A method of manufacturing structural steel having excellent toughness in the heat-affected zone, comprising: hot-rolling at a ratio and then cooling at a cooling rate of 10°C/sec or less from the rolling end temperature to room temperature.
According to claim 4, The steel slab, Ni: 0.1 ~ 3.0%, Cu: 0.1 ~ 1.5%, Mo: 0.05 ~ 1.0%, Cr: 0.05 ~ 1.0%, W: 0.001 ~ 0.2%, Nb: 0.005- Welding characterized in that one or more selected from the group consisting of 0.1%, V: 0.005 to 0.1%, and at least one selected from Ca: 0.0005 to 0.005% and REM: 0.005 to 0.05% are additionally included A method of manufacturing structural steel with excellent toughness in the heat affected zone.
High heat input welding is applied to the steel material (base material) manufactured by the manufacturing method of claim 3 or 4 to generate a prior austenite of 120 μm or less in the welding heat-affected zone, followed by rapid cooling to produce the welding heat-affected zone. A welded structure in which ferrite with a microstructure of 20㎛ or less has a phase fraction of 70% or more.
The welded structure according to claim 6, wherein the ferrite is at least one of polygonal ferrite and acicular ferrite structures.