KR100362680B1 - High strength steel plate having superior toughness in weld heat-affected zone and Method for manufacturing the same, welding fabric using the same - Google Patents

Abstract

The present invention relates to structural steels used in welding structures, such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes, etc. The purpose of the base material is bainite + ferrite TiN precipitates excellent in high temperature stability while ensuring high strength It is to provide a steel and a method for manufacturing the welded structure while improving the toughness of the weld heat affected zone by finely and uniformly distributed.

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: 0.0005-0.1%, N: 0.008 to 0.03%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less, 1.2≤Ti / N≤2.5, 10≤N / B Satisfies ≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B + V) / N≤14, is composed of the remaining Fe and other unavoidable impurities, and the microstructure is 30-80% bainite The technical gist of the welded structural steel having excellent weld heat affected zone toughness composed of a composite structure of ferrite of less than 20 μm and the rest thereof is considered as the technical gist.

Description

High strength steel plate having superior toughness in weld heat-affected zone and Method for manufacturing the same, welding fabric using the same}

The present invention relates to structural steel used in welded structures, such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes. More specifically, the base material is made of bainite + ferrite, and relates to a welded structural steel and a method of manufacturing the same, while ensuring high strength and distributing TiN precipitates having excellent high temperature stability finely and uniformly to improve the toughness of the weld heat affected zone.

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

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. That is, if the welding heat input can be increased, the range of use thereof can be widened. The range of high heat input currently used corresponds to approximately 100-200 kJ / cm, and in order to weld steel with a thicker plate thickness of 50 mm or more, it is possible to have a super heat input range of 200-500 kJ / cm.

When the heat input is applied to the steel, the heat affected zone formed during welding, particularly the heat affected zone near the fusion boundary, is heated to a temperature close to the melting point by the amount of heat input. 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 becomes 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 growth of the austenite grains in the weld heat affected zone and to keep it fine. As a means to solve this problem, 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 high temperature, in steel is used. Such techniques include Japanese Patent Laid-Open No. 11-140582, No. 10-298708, No. 10-298706, No. 9-194990, No. 9-324238, No. 8-60292 (S) 60-245768, (Pyeong) 5-186848, (S) 58-31065, (S) 61-79745, Journal of the Japan Welding Society, Vol. 52, No. 2, 49 and Japanese Patent Laid-Open 64-15320, etc.

Japanese Patent Laid-Open No. 11-140582 is a representative technique using TiN precipitates. When 100 J / cm of heat input (maximum heating temperature of 1400 ° C) is applied, impact toughness at 200 ° C is about 200J. A structural steel material of about 300J) is disclosed. In this 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 pieces / mm 2 to 8.1 × 10 ⁴ pieces / mm 2, with TiN precipitates of 0.03 to 0.2 μm being 3.9. × 10 ³ gae /㎟~6.2×10 by precipitation in ^four / ㎟ and by refining the ferrite to secure the toughness of the weld. This steel is a microstructure of ferrite and pearlite, and has a mechanical property of tensile strength up to 581 MPa and yield strength up to 405 MPa.

However, according to this 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 at the highest impact toughness of 0 ° C), As the difference between the base material 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, in order to secure the precipitate of TiN, the slab is heated at a temperature of more than 1050 ℃ quenched, and then adopts a process of reheating for hot rolling, there is a disadvantage that the manufacturing cost is high due to two heat treatment.

To date, many techniques for improving the toughness of the weld heat affected zone during high heat input welding have been known. There is no bar. In particular, there is no technology in which the base material is high in strength and the toughness of the weld column is equivalent to that of the base material. Therefore, if the above problems can be solved, super heat input welding of thickened high strength steels is possible, thereby achieving high welding operation as well as securing a high layer of steel structures and securing reliability of steel structures.

An object of the present invention is to provide a welded structural steel and a method of manufacturing the same, and a welded structure using the steel, which has a high strength of the steel (base metal) and excellent toughness of the weld heat affected zone.

Welded structural steel materials of the present invention for achieving the above object, by weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1 %, N: 0.008 to 0.03%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less, 1.2≤Ti / N≤2.5, 10 Satisfies ≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B + V) / N≤14, is composed of the remaining Fe and other unavoidable impurities, and the microstructure is 30-80 It consists of a composite structure of% bainite and the remaining ferrite of 20 µm or less.

In addition, the manufacturing method of the welded structural steel material of this invention is C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1% by weight. , N: 0.008-0.03%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less, 1.2≤Ti / N≤2.5, 10≤ A steel slab satisfying N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B + V) / N≤14 and composed of the remaining Fe and other unavoidable impurities is 1150-1250 ° C. After heating for 60 to 180 minutes at the temperature, hot rolling at an austenite recrystallization zone with a rolling ratio of 40% or more, and then cooling to a bainite transformation end temperature ℃ ± 10 ℃ at a rate of 5-20 ℃ / sec. do.

In addition, the welded structure of the present invention is welded to a base material having a composition and a microstructure of the steel of the present invention, in which TiN precipitates of 0.01-0.1 μm are distributed at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less. It is applied to produce austenite (prior austenite) of less than 80㎛ in the weld heat affected zone, and then quenched so that the microstructure of the weld heat affected zone is less than 20㎛ ferrite of 70% or more.

Hereinafter, the present invention will be described in detail.

In the present invention, the term "prior Austenite" refers to austenite formed in the weld heat affected zone when welding is applied to the steel, and the austenite formed during the manufacturing process of the steel (hot rolling process). Use for convenience to distinguish.

The inventors of the present invention have studied the method of increasing the strength of the steel (base material) and the high toughness of the welded heat affected zone at the same time. As a result, when the microstructure of the base material is a composite structure of bainite and ferrite, it is very effective in improving the base material strength. In fact, it was confirmed that the toughness of the weld heat affected zone does not become a problem if the size of the austenite grains of the weld heat affected zone is controlled below the threshold (about 80 μm) when welding to the base material having such a microstructure.

Starting from this point of view, the inventors of the present invention have been able to derive the following method of improving the toughness of the weld heat affected zone by managing the size of the austenite grains below the critical value in the composite structure of bainite + ferrite.

[1] While minimizing the distribution of fine TiN precipitates, the solubility product exhibiting high temperature stability of these precipitates is reduced,

[2] The grain size of the old austenite is refined by setting the initial ferrite grain size of the steel (base material) below the critical level.

[3] According to the present invention, when BN and AlN precipitates are effectively used while managing the crystal grain size of the old austenite below the threshold value, the ferrite fraction in the weld heat affected zone, in particular, the polygonal or acicular ferrite is effective. It is possible to improve the toughness of the weld heat affected zone.

[4] The base material strength is improved by securing the proper fraction of bainite microstructure in the base material through accelerated cooling in the rolling process. These [1] [2] [3] [4] are demonstrated in more detail.

[1] TiN precipitates

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 more, and TiN precipitates precipitated in the base metal are partially dissolved by welding heat or Ostwald ripening. Small precipitates are decomposed and diffuse into larger precipitates, resulting in larger precipitates). Only some precipitates are coarse, and the number of TiN precipitates is significantly reduced, thereby suppressing the inhibitory effect of the growth of austenite grains. .

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 characteristics of TiN precipitates according to the ratio of Ti / N are as follows. / N ratio is low), it is found that the dissolved Ti concentration and diffusion rate of the dissolved Ti atoms and the high temperature stability of TiN precipitates are improved. That is, when the ratio of Ti and N (Ti / N) is in the range of 1.2 to 2.5, the amount of solid solution Ti is extremely reduced and the high temperature stability of the TiN precipitate is increased, so that the fine TiN precipitate of 0.01-0.1 μm or less is 0.5 μm or less. A result of distribution of 1.0 × 10 7 pieces / mm 2 or more at intervals was obtained. Increasing the nitrogen content at the same Ti content, all of the Ti atoms dissolved are easily combined with the nitrogen atom, and in the high nitrogen environment, the amount of solid solution Ti decreases, so that the TiN precipitates are stabilized at a higher temperature than when the nitrogen content is low. It is analyzed that this is because the solubility product is lowered. In view of the fact that the present invention can promote the aging due to the presence of solid solution N in a high nitrogen environment, the ratio of N / B, Al / N, V / N, and N and Ti + Al + B + (V) Is managed as a whole to precipitate N into BN, AlN, and VN.

As described above, in the present invention, by controlling the distribution of TiN precipitates and the solubility of TiN according to the ratio of Ti / N, the toughness difference between the base material and the heat affected zone is lowered to within 30 J, which is simply increased by increasing the content of Ti (Ti / N≥ 4) There is a significant difference from the conventional precipitate management technique (Japanese Laid-Open Patent Publication No. 11-140582) which increases the amount of TiN precipitates.

[2] microstructures, steel (base material)

According to the study of the present invention, in order to make the size of the old austenite 80㎛, it can be seen that it is important to make the size of the ferrite to 20㎛ or less while the microstructure of the base material as a composite structure of ferrite + bainite. At this time, the refinement of the ferrite is to suppress the growth of the ferrite grains generated in the cooling process by using carbides (WC, VC) as well as austenite grain refinement due to steel working during hot rolling.

[3] microstructure of weld heat affected zone

The fact of the present invention reveals that the toughness of the weld heat affected zone is important not only in the size of the austenite grains when the base material is heated above 1400 ° C, but also in the amount and size of the ferrite precipitated at the former austenite grain boundaries and its shape. It affects. That is, in consideration of the toughness of the weld heat affected zone, it is important that the amount of ferrite is 70% or more and the size is 20 μm or less. In particular, it is important to induce the transformation of polygonal ferrite and acicular ferrite in the austenite mouth. In the present invention, AlN, Fe ₂₃ (B, C) ₆ precipitates, BN precipitates are used for this purpose.

[4] bainite tissue fraction control

The present inventors can easily control the bainite structure fraction which can improve the strength of the base metal when the accelerated cooling rate is controlled (5-20 ° C./sec) in the hot rolling process, and the properties of the weld heat affected zone are It was confirmed that it is not related to the microstructure change of.

Hereinafter, the present invention will be described in detail by dividing the steel and its manufacturing method.

[Welding Structural Steels]

It is preferable to make content of carbon (C) into 0.03 to 0.17%.

If the content of C is less than 0.03%, securing strength as a structural steel is insufficient. In addition, when C exceeds 0.17%, microstructures susceptible to toughness such as upper bainite, martensite and degenerate pearlite during transformation are transformed to lower the low temperature impact toughness of structural steel, and Increasing hardness or strength results in toughness degradation and generation of weld cracks.

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

If the Si content is less than 0.01%, the deoxidation effect of the molten steel is insufficient during steelmaking and the corrosion resistance of the steel is deteriorated. If the content of Si is more than 0.5%, the effect is saturated, and the degree of quenchability during cooling after rolling is increased. It promotes the transformation of martensite and lowers the low temperature impact toughness.

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

Mn is an effective element which deoxidizes in steel and improves weldability, hot workability, and strength. Mn forms a substituted solid solution in the matrix to strengthen the matrix to secure the strength and toughness. For this purpose, Mn is preferably added at least 0.4%. However, when the Mn content is more than 2.0%, it has a detrimental effect on the toughness of the weld heat affected zone due to tissue heterogeneity due to manganese segregation rather than a solid solution strengthening effect. In addition, when the steel solidifies, macro segregation and micro segregation occur depending on the segregation mechanism, which promotes the formation of a central segregation zone in the center of rolling, thereby causing a low temperature transformation structure in the center of the base metal. In particular, manganese is an element that precipitates in the form of MnS around the Ti-based oxide and affects the formation of acicular and polygonal ferrites effective for improving the toughness of the weld heat affected zone.

The content of aluminum (Al) is preferably limited to 0.0005-0.1%.

Al is not only an element necessary for deoxidation, but also an indispensable element for forming fine AlN precipitates in steel. In addition, Al is an element which reacts with oxygen to form an Al oxide, and thus is necessary for Ti to form fine TiN precipitates without reacting with oxygen. In order to form a fine AlN precipitate, the Al should be added 0.0005% or more, but if the content exceeds 0.1%, Weedman Stetten is vulnerable to toughness during the cooling of the welding heat affected zone of AlN precipitated and the remaining solid Al It promotes the production of ferrite (Widmanstatten ferrite) and phase martensite to lower the toughness of the high heat input welding heat affected zone.

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

Ti is indispensable 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 precipitation effect, Ti should be added more than 0.005%, but when the content exceeds 0.2%, coarse TiN precipitate and Ti oxide are formed in molten steel, which does not inhibit austenite grain growth of the weld heat affected zone. It is not preferable because it is.

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

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%, it is not preferable because the hardenability increases due to hardening and low temperature cracking of the weld heat affected zone.

The content of nitrogen (N) is preferably limited to 0.008-0.03%.

N is an indispensable element for forming TiN, AlN, BN, VN, NbN, etc., and it suppresses austenite grain growth of welding heat affected zone at the time of high heat input welding and precipitates such as TiN, AlN, BN, VN, NbN Increase the amount. In particular, since the TiN and AlN precipitates have a remarkable effect on the size, precipitate interval, precipitate distribution, complex precipitation frequency with oxide, and high temperature stability of the precipitate itself, the content is preferably set to 0.008% or more. However, when the N content exceeds 0.03%, the effect is saturated, toughness is reduced 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 to reduce the toughness of the weld metal. Cause.

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

W is an element that uniformly precipitates in the base material as tungsten carbide (WC) after hot rolling to effectively suppress ferrite grain growth after ferrite transformation, and also suppress growth of austenite grains in the initial heating of the weld heat affected zone. If the content is less than 0.001%, the tungsten carbide for inhibiting ferrite grain growth is less distributed during the hot rolling and it is not preferable because the effect is saturated when more than 0.2% is added.

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

P is preferably as low as possible because it is an impurity element that promotes central segregation during rolling and hot cracking during welding. 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.

Since S forms a low melting point compound such as FeS when present in a large amount, it is preferable to manage S as low as possible. 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. However, in the case of sulfur, it precipitates in the form of MnS around the Ti-based oxide and thus affects the formation of needle-shaped and polygonal ferrites effective for improving the toughness of the weld heat affected zone. It is desirable to limit the amount from 0.003% to 0.03% or less.

Oxygen (O) is preferably at most 0.005%.

When the oxygen content exceeds 0.005%, the Ti element is not preferable because the Ti element is formed of Ti oxide in molten steel and thus does not form a TiN precipitate. Also, inclusions such as coarse Fe oxide and Al oxide are formed, which is bad for the base metal. It is not desirable because it affects.

The ratio of Ti / N is preferably 1.2 to 2.5.

In the present invention, the Ti / N ratio is lowered to 2.5 or less, which has two advantages. First, it is possible to distribute uniformly while increasing the number of TiN precipitates. In other words, if the nitrogen content is increased at the same Ti content, all Ti atoms that are employed in the cooling process during the playing process are combined with the nitrogen atoms to distribute fine TiN precipitates at narrow intervals. Secondly, the solubility product showing stability at high temperature is small, thereby preventing re-use of the Ti. That is, in a high nitrogen environment, Ti has a stronger tendency to bind with N rather than solid solution, so that the TiN precipitate becomes stable at high temperature. However, when the Ti / N ratio is less than 1.2, it is not preferable because the amount of solid solution nitrogen in the base material increases, which is detrimental to the toughness of the weld heat-directed portion. On the other hand, if the Ti / N ratio is higher than 2.5, coarse TiN is crystallized in molten steel, which is a steelmaking process, and a uniform distribution of TiN is not obtained. It is not preferable because it adversely affects the impact toughness.

It is preferable to make ratio of N / B into 10-40.

In the present invention, if the N / B ratio is less than 10, the precipitation amount of BN which promotes the ferrite transformation of polygons at the austenite grain boundary during the cooling process after welding is insufficient, and when the N / B ratio exceeds 40, the effect is saturated and dissolved. This is because the amount of nitrogen is increased to lower the toughness of the weld heat affected zone.

It is preferable to make Al / N ratio into 2.5-7.

In the present invention, when the Al / N ratio is less than 2.5, AlN precipitates for inducing needle-like ferrite transformation are insufficient, and the amount of solid solution nitrogen in the weld heat affected zone may increase, resulting in a weld crack, and an Al / N ratio of more than 7 In the case the effect is saturated.

It is preferable that ratio of (Ti + 2Al + 4B) / N is 6.5-14.

In the present invention, when the ratio of (Ti + 2Al + 4B) / N is less than 6.5, the growth inhibition of the austenite grain growth of the weld heat affected zone, the generation of fine polygonal ferrite at the grain boundary, the amount of solid solution nitrogen, the needle shape in the grain and the polygonal ferrite The size and number of distributions of TiN, AlN, BN, and VN precipitates for formation and control of the tissue fraction are insufficient, and the effect is saturated when (Ti + 2Al + 4B) / N exceeds 14. On the other hand, when V is added, it is preferable to set the ratio of (Ti + 2Al + 4B + V) / N to 7-17.

In order to further improve the toughness of the steel material (base material) and the heat-affected portion formed as described above, V is further added.

The content of vanadium (V) is preferably limited to 0.01-0.2%.

V is an element that combines with N to form VN to promote ferrite formation in the weld heat affected zone, and VN precipitates alone or precipitates in TIN precipitates to promote ferrite transformation. In addition, V combines with C to form VC, which acts to inhibit ferrite grain growth after ferrite transformation. When the V content is less than 0.01%, it is difficult to obtain the ferrite transformation promoting effect in the weld heat affected zone because the VN deposition amount is small. On the other hand, exceeding 0.2% is not preferable because it causes toughness of the base metal and the weld heat affected zone (HAZ) and improves the weld hardenability, which may cause the low temperature crack of the weld.

Moreover, it is preferable to make ratio of V / N into 0.3-9.

In the present invention, when the V / N ratio is less than 0.3, it is difficult to secure an appropriate number and size of VN precipitates deposited and distributed at the TiN + MnS precipitate boundary for improving the toughness of the weld heat affected zone. On the other hand, when the V / N ratio exceeds 9, the weld heat influences because the size of the VN precipitates deposited at the TiN + MnS precipitate boundary is coarsened, and the number of VN precipitations deposited at the TiN + MnS composite precipitate boundary is reduced. Reduce the ferrite phase fraction effective for negative toughness.

In the present invention, in order to further improve the mechanical properties in the steel composition as described above, one or more selected from the group of Ni, Cu, Nb, Mo, Cr is further added.

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

Ni is an effective element for improving the strength and toughness of the base material by solid solution strengthening. To achieve this effect, Ni should be added more than 0.1%, but if the content exceeds 3.0%, the hardenability increases to decrease the toughness of the weld heat affected zone, and there is a possibility of high temperature crack in the weld heat affected zone and the weld metal. Because it is not desirable.

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

Cu is an element that is effective at securing the strength and toughness of the base material due to solid solution at the base. In order to achieve this effect, Cu should be added more than 0.1%, but if the content exceeds 1.5%, it is desirable because it increases the hardenability in the weld heat affected zone and lowers the toughness and promotes high temperature crack in the weld heat affected zone and the weld metal. I can't. In particular, the Cu precipitates in the form of CuS around the Ti-based oxide together with sulfur, thereby affecting the formation of needle-shaped and polygonal ferrites, which are effective for improving the toughness of the weld heat affected zone, and thus the content is 0.3-1.5%. desirable.

In addition, when adding Cu and Ni compositely, it is preferable to make these sum total less than 3.5%. The reason for this is that when it exceeds 3.5%, the hardenability becomes large, which adversely affects the weld heat affected zone toughness and weldability.

The content of niobium (Nb) is preferably limited to 0.01-0.10%.

Nb is an effective element from the viewpoint of securing the base material strength, and this effect cannot be obtained when the Nb content is less than 0.01%. On the other hand, if it exceeds 0.1%, it causes undesired precipitation of coarse NbC, which is detrimental to the toughness of the base metal.

Chromium (Cr) is preferably made 0.05 to 1.0%.

Cr increases the hardenability and also improves the strength, when the content is less than 0.05%, strength cannot be obtained. If it exceeds 1.0%, the base metal and HAZ toughness deteriorate.

Molybdenum (Mo) is preferably 0.05-1.0%.

Mo is an element that has an effect of increasing hardenability and improving strength. The content of Mo is preferably set to 0.05% or more for securing strength, but in order to suppress HAZ hardening and welding low temperature cracking, the upper limit thereof is 1.0%. It is preferable to set it as.

In addition, in the present invention, one or two kinds of Ca and REM are further added to suppress grain growth of austenite during heating.

Ca and REM form an oxide having excellent high temperature stability, thereby suppressing austenite grain growth when heated in the base metal and improving 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, but when Ca exceeds 0.005% of Rem of more than 0.05%, large inclusions and clusters are generated to harm the cleanliness of the steel. As REM, 1 type, or 2 or more types, such as Ce, La, Y, and Hf, may be used, and any of the above effects can be obtained.

· Microstructure of steel

In the present invention, the steel is a composite structure of ferrite + bainite, the tissue fraction of bainite is preferably in the range of 30-80%. The reason is that less than 30% is difficult to secure the appropriate base material strength for showing the effect of the present invention, when the base material toughness is more than 80%.

In the composite structure of ferrite + bainite, the ferrite grain size is preferably 20 µm or less. This is because when the grain size of the ferrite is larger than 20 μm, the austenite grain size of the weld heat affected zone becomes 80 μm or more during high heat input welding, which is detrimental to the weld heat affected zone toughness.

Distribution of precipitates

The former austenite grains of the weld heat affected zone are greatly influenced by the size and number of nitrides distributed in the base metal and their distribution when the austenite grain size of the base material is constant. In addition, 30-40% of the nitrides distributed in the base material when the heat input welding is higher than the heat input temperature (above the heating temperature of 1400 ℃) are re-used as the base material, thereby reducing the austenite grain growth inhibiting effect of the welding heat affected zone. There is a need for a uniform distribution of further nitrides taking into account the nitrides that become. In order to suppress the growth of austenite in the weld heat affected zone, it is important to uniformly distribute the fine TiN precipitate to suppress the Ostwald ripening phenomenon in which some precipitates are coarsened. For this purpose, the TiN precipitates should be controlled to 0.5 μm or less to make the TiN distribution uniform.

In addition, it is preferable to limit the particle diameter and the critical number of Ti N to 0.01-0.1 μm and 1.0 × 10 ⁷ or more per 1 mm ² . The reason for this is that less than 0.01 μm is easily re-used to the base metal during the high heat input welding, so that the effect of inhibiting the growth of the austenite grains is insufficient. If the thickness exceeds 0.1 μm, pinning of the austenite grains occurs. This is because the 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 is difficult to control the size of the austenite grains of the weld heat-affected zone at the time of welding higher than the heat input to be less than or equal to the threshold of 80 μm.

[Method of manufacturing welded structural steel]

Refining (Deoxidation, Degassing) Process

In general, the steel refining process consists of an out-of-furnace refining process after the first refining of the converter and the second refining of the molten steel of the converter by ladle. ). Usually deoxidation takes place between primary and secondary refining.

In the present invention, the dissolved oxygen is adjusted to below an appropriate level in such a deoxidation process, and then Ti is added so that most of the solution is dissolved in molten steel without forming Ti as an oxide. For this purpose, it is preferable to inject and deoxidize an element having a greater deoxidizing power than Ti before adding Ti. The deoxidizing power of the deoxidizer is as follows.

Cr <Mn <Si <Ti <Al <REM <Zr <Ca ≒ Mg

The amount of dissolved oxygen is greatly influenced by the formation behavior of the oxide, and the deoxidizer having a high affinity with oxygen has a very high rate of bonding with oxygen in the molten steel. Therefore, if deoxidation is performed using an element having a greater deoxidizing power before adding Ti, it is possible to prevent Ti from forming an oxide as much as possible. Of course, before the addition of the element (Al), which has a greater deoxidizing power than Ti, by adding and deoxidizing Mn, Si and the like, which are the five major elements of steel, and then deoxidizing by adding Al, the amount of deoxidizer added is preferable.

On the other hand, floating separation of inclusions in molten steel is generally known to proceed in the following order. (Dissolution of deoxygenation element in molten steel) → (Involvement of nucleation) → (Growth of inclusions) → (Continuous growth and injury due to collisions between inclusions) → (Removal of slag from slag on molten steel surface) Since the speed of each step varies depending on the type, deoxidation using a strong deoxidation element can lower the dissolved oxygen amount more easily.

In the present invention, the amount of dissolved oxygen is added as low as possible by injecting a strong deoxidation element before the introduction of Ti, and in order to maximize the amount of Ti dissolved in molten steel, it is preferable to set it as at least 30 ppm or less. The reason for this is that when the dissolved oxygen amount exceeds 30 ppm, the oxygen in the molten steel and Ti are combined to form Ti oxide easily when Ti is added, and the amount of solid solution Ti decreases.

On the other hand, according to the 'Stoke Law' widely used in steelmaking, the higher the density of inclusions, the more difficult the inclusions are. Injuries are difficult As the inclusions increase in the steel, the addition of a deoxidation element forming a high density inclusion may be used as an additional advantage according to the distribution of oxides. Has no effect on the effect.

After adjusting the amount of dissolved oxygen according to the present invention, it is preferable to add Ti within 10 minutes so that the content is 0.005-0.2%. If less than 0.005% of Ti is contained in the molten steel after deoxidation, it is difficult to form a large amount of fine TiN in the casting process, and if it contains more than 0.2%, the effect is saturated and the TiN is coarsened to expect the austenite grain suppression effect. Difficult to do The reason for adding Ti within 10 minutes is that Ti oxide is generated and the amount of solid solution Ti decreases as the time after Ti is added. In the case where the vacuum degassing treatment ('RH') is performed in the refining process, the addition of Ti can be performed either before or after the vacuum degassing treatment.

Casting process

In the present invention, the molten steel refined as described above is continuously cast into slabs. Continuous casting is preferable from the viewpoint of productivity improvement by casting at low speed and giving a weak cooling condition in the secondary cooling zone in consideration of the high possibility of occurrence of cast surface cracks in high nitrogen steel. Cooling conditions in the secondary cooling zone are important factors affecting the refinement and uniform distribution of TiN precipitates.

According to the study of the present invention, the continuous casting speed is less than 1.1 m / min, more preferably about 0.9 to 1.1 m / min, which is lower than the usual casting speed of about 1.2 m / min. The reason is that when the casting speed is less than 0.9m / min, the surface cracks are advantageous, but the productivity is lowered.

In the secondary cooling zone, the specific water amount is preferably as low as possible, that is, 0.3 to 0.35 l / kg. When the specific amount is less than 0.3 L / kg, it is difficult to control the proper size and number of TiN for showing the effect of the present invention by coarsening TiN precipitates. In addition, when the specific water content exceeds 0.35L / kg, the precipitation frequency of the TiN precipitates is small, and it is difficult to control the number, size, and the like of TiN precipitates for showing the effects of the present invention.

Hot rolling process

In the present invention, the slab is heated at 1100-1250 ° C. for 60-180 minutes. It is a problem because the number of TiN precipitates is small because the rate of diffusion of solute atoms is less than 1100 ℃, Ti precipitates, etc. are coarse or decomposed when it exceeds 1250 ℃, the precipitates decrease the number of precipitates Not desirable On the other hand, if the heating time is less than 60 minutes, it is not preferable because there is no segregation reduction effect of the solute atoms and there is not enough time for the solute atoms to diffuse to form precipitates. In addition, when the heating time exceeds 180 minutes, coarsening of austenite grain size occurs, which is not preferable in terms of work productivity.

After heating as above, it is preferable to hot-roll at a rolling ratio of 40% or more at the austenite recrystallization zone temperature. The austenite recrystallization zone temperature is affected by the emphasis and the previous reduction amount, and the austenite recrystallization zone temperature is in the range of about 1050 to 850 ° C. in consideration of the usual reduction amount in the emphasis of the present invention. In this section, a rolling ratio of at least 40% should be given. If the rolling ratio is less than 40%, the ferrite nucleation site in the austenite grain is insufficient and the effect of refining the ferrite grains due to austenite recrystallization is insufficient. It affects the precipitate behavior which effectively affects the toughness of the affected zone.

In hot rolling, the austenite grain size is affected by the temperature, time in the reheating furnace, and the rolling amount. The grain size of the austenite affects the quenchability, so that the desired bainite fraction can be obtained by controlling it. In order to increase the bainite fraction, it is recommended to set the austenite grain size to 10 μm or more.However, if the austenite grain size becomes larger than 50 μm, the hardenability becomes too large during transformation, which may cause martensite transformation. .

After hot rolling in the present invention, the reason for limiting the cooling rate in the range of 5-20 ° C / sec to bainite transformation end temperature ± 10 ° C is as follows. The phase transformation of the present invention steel occurs in a section within the bainite transformation end temperature ± 10 ℃, the cooling rate must be controlled up to this section. If the accelerated cooling rate is less than 5 ° C / sec it is difficult to secure the bainite phase fraction for showing the effect of the present invention, and in the case of more than 20 ° C / sec martensite phase ratio increases to be harmful to the base material toughness.

Casting of the steel in the present invention can be produced by slab by continuous casting or mold casting. In this case, if the cooling rate is fast, it is advantageous to finely disperse the precipitates, and thus, continuous casting having a high cooling rate is preferable. For the same reason, the slab is advantageously thinner. In addition, in the hot rolling process, hot charge rolling and direct rolling may be applied according to a user's use, and various known technologies such as control rolling and control cooling may be applied. In addition, heat treatment may be applied to improve the mechanical properties of the hot rolled sheet produced according to the present invention. However, even if the well-known techniques are applied to the present invention, it is natural that they are interpreted to be substantially within the technical scope of the present invention as a simple change of the present invention.

[Welding Structure]

The welded structural steel provided according to the present invention is a ferrite composite structure having a grain size of 30 to 80% bainite + 20 µm or less, and a precipitate of TiN is 1.0x10 ⁷ per 1 mm ² having a size of 0.01-0.1 µm. It is ideally distributed and its interval is 0.5 탆 or less.

When welding is applied to such steels, the grain size of the old austenite becomes 80 µm or less in the weld heat affected zone. When the grain size of the former austenite is 80 μm or more, it is easy to form and use low temperature structure (martensite or upper bainite) due to the increase in hardenability, which is harmful to the toughness of the weld heat affected zone and also to the austenite grain boundary. Even if ferrites with different nucleation sites are produced at, ferrites coalesce on grain growth and have a detrimental effect on toughness. Therefore, the critical size of the austenite grain size of the weld heat affected zone for showing the effect of the present invention must be controlled to 80㎛ or less.

When the high heat input welding is applied and quenched as described above, the microstructure of the heat-affected portion has a phase ratio of 70% or more of ferrite having a size of 20 μm or less. When the grain size of the ferrite is larger than 20 μm, the ferrite fraction of the side plate or allotriomorphs harmful to the weld heat affected zone toughness increases. In addition, in order to improve toughness, the ferrite phase ratio is preferably 70% or more. The ferrite of the present invention is more advantageous for toughness when it has the characteristics of polygonal ferrite and acicular ferrite. This can be induced by forming BN, Fe carbide boride in the grain boundary and in the mouth according to the present invention.

Hereinafter, the present invention will be described in more detail with reference to Examples.

EXAMPLE

In order to make the steel slab with the composition of the composition as shown in Table 1, except for the component of Ti in Table 1, other components were dissolved in the converter using a sample of the same invention steel, then decanted from Mn to Si under the conditions of Table 2, Then, deoxidized with Al to adjust the dissolved oxygen amount, and then added Ti to adjust the concentration of Ti as shown in Table 1, maintaining molten steel for a certain time, casting it to an ingot, and adjusting the cooling rate to produce slabs, It was prepared as a hot rolled sheet under the conditions of Table 4.

The test pieces for evaluating the mechanical properties of the base metal from the hot rolled plates as described above were taken at the center of the plate thickness of the rolled material, the tensile test piece was taken in the rolling direction, and the Charpy impact piece was taken in the direction perpendicular to the rolling direction. It was.

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, 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 heating up to the maximum heating temperature (1200 ~ 1400 ℃) by using the simulation welding simulator (simulator). It was then quenched using He gas. The quenched specimens were polished and corroded to determine the austenite grain size at the highest heating temperature condition by KS (KS D 0205).

TiN precipitate size, number, and spacing, 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 applied, the surface of the test piece was polished and processed into an impact test piece, and evaluated through a Charpy impact test at -40 ° C.

Alloy element composition ratio for showing the effect of the present invention Ti / N N / B Al / N V / N (Ti + 2Al + 4B + V) / N Inventive Steel 1 1.2 17.1 3.3 0.8 8.9 Inventive Steel 2 1.8 28.0 2.5 0.4 7.3 Invention Steel 3 1.4 36.7 5.5 1.8 14.2 Inventive Steel 4 2.5 16.0 2.5 6.3 14.0 Inventive Steel 5 1.7 20.0 3.0 1.7 9.5 Inventive Steel 6 2.0 10.0 2.5 9.0 16.4 Inventive Steel 7 1.3 14.4 3.5 1.7 10.3 Inventive Steel 8 1.5 12.0 5.0 0.8 12.7 Inventive Steel 9 2.2 22.5 2.8 2.2 10.2 Inventive Steel 10 2.5 16.7 4.5 2.0 13.7 Inventive Steel 11 1.4 12.0 3.8 - 9.3 Conventional Steel 1 4.1 13.8 0.6 - 5.7 Conventional Steel 2 2.5 96.0 0.8 - 4.0 Conventional Steel 3 0.8 105.8 0.4 - 1.5 Conventional Steel 4 4.1 4.0 0.8 8.8 15.5 Conventional Steel 5 6.5 4.0 1.1 18.5 28.1 Conventional Steel 6 3.2 2.6 0.4 16.1 21.6 Conventional Steel 7 1.0 9.9 2.5 - 6.5 Conventional Steel 8 1.2 14.3 0.4 - 2.2 Conventional Steel 9 0.8 9.1 2.1 3.9 9.2 Conventional Steel 10 0.6 9.5 3.2 1.5 8.9 Conventional Steel 11 5.5 12.7 3.4 7.8 20.3

Steel grade used division Primary deoxidation sequence Dissolved oxygen content after addition of secondary deacidification (ppm) Ti addition after finishing deoxidation (%) Casting speed (m / min) Specific water volume in the secondary cooling stand (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 22 0.016 0.95 0.35 Inventive Steel 2 Invention 2 Mn → Si 21 0.053 0.98 0.32 Invention Steel 3 Invention 3 Mn → Si 20 0.016 1.02 0.32 Inventive Steel 4 Invention 4 Mn → Si 23 0.022 0.98 0.35 Inventive Steel 5 Invention 5 Mn → Si 21 0.053 1.04 0.32 Inventive Steel 6 Invention 6 Mn → Si 24 0.021 1.05 0.35 Inventive Steel 7 Invention 7 Mn → Si 21 0.017 1.05 0.38 Inventive Steel 8 Invention Material 8 Mn → Si 22 0.019 1.04 0.34 Inventive Steel 9 Invention Material 9 Mn → Si 18 0.022 0.99 0.35 Inventive Steel 10 Invention 10 Mn → Si 21 0.026 1.04 0.35 Inventive Steel 11 Invention 11 Mn → Si 19 0.021 1.02 0.32 Refining conditions of conventional steels are not specifically described in the specification.

Steel grade used division Slab heating temperature (℃) Slab heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount (%) Rolling amount at recrystallization station (%) Cooling rate (℃ / sec) Cooling end temperature (℃) Invention 1 Inventive Example 1 1150 160 1010 780 85 55 7 450 Inventive Example 2 1200 130 1000 790 85 55 15 450 Inventive Example 3 1250 100 1120 780 85 55 20 440 Comparative Example 1 1000 60 950 780 85 55 15 450 Comparative Example 2 1400 350 1200 860 85 55 16 440 Invention 2 Inventive Example 4 1210 130 1020 810 80 50 16 450 Invention 3 Inventive Example 5 1220 130 1020 800 80 50 17 450 Comparative Example 3 1200 120 1010 810 80 50 One Room temperature Comparative Example 4 1210 130 1010 800 80 45 35 Room temperature Invention 4 Inventive Example 6 1190 140 1010 820 80 50 15 430 Invention 5 Inventive Example 7 1180 150 1030 780 80 50 15 450 Invention 6 Inventive Example 8 1185 150 1000 780 80 50 18 430 Invention 7 Inventive Example 9 1215 120 1020 790 80 50 16 430 Invention Material 8 Inventive Example 10 1200 120 1010 780 75 45 12 430 Invention Material 9 Inventive Example 11 1210 120 1020 820 75 45 14 430 Invention 10 Inventive Example 12 1210 110 1000 800 80 50 16 450 Invention 11 Inventive Example 13 1220 110 1010 810 75 45 14 450 Conventional Steel 11 1200 - Ar ₃ or higher 960 80 Cooling Manufacturing conditions of conventional steel (1-10) are not specifically presented

division Base material precipitate characteristics Base material mechanical and tissue properties Number of precipitates (pcs / mm ² ) Precipitate size (nm) Average precipitate interval (㎛) Thickness (mm) Yield strength (MPa) Tensile Strength (MPa) Elongation (%) Bainite fraction (%) -40 ℃ impact toughness (J) Inventive Example 1 2.6 X ^{10 8} 17 0.25 25 582 752 21.4 38 252 Inventive Example 2 3.2 X ^{10 8} 13 0.35 25 596 768 22.2 38 234 Inventive Example 3 2.4 X ^{10 8} 12 0.35 25 586 761 20.4 39 225 Comparative Example 1 2.2 X 10 ⁶ 117 1.37 25 592 768 35.0 10 228 Comparative Example 2 2.3 X 10 ⁵ 158 2.75 25 585 772 12.3 15 236 Inventive Example 4 3.4 X ^{10 8} 18 0.32 30 596 768 20.2 40.2 220 Inventive Example 5 2.6 X ^{10 8} 17 0.30 30 586 769 19.8 43.0 215 Comparative Example 3 3.2 X ^{10 8} 18 0.37 30 396 468 35.0 5 328 Comparative Example 4 3.2 X ^{10 8} 19 0.36 30 678 872 16.5 14.0 98 Inventive Example 6 3.3 X ^{10 8} 15 0.28 30 582 771 20.4 42.6 262 Inventive Example 7 4.5 X ^{10 8} 21 0.26 35 587 782 21.2 40.5 228 Inventive Example 8 4.3X10 ⁸ 17 0.31 35 612 768 18.9 42.1 235 Inventive Example 9 5.4 X ^{10 8} 16 0.30 35 599 785 18.5 43.1 220 Inventive Example 10 5.5X10 ⁸ 20 0.25 35 624 783 19.5 42.0 236 Inventive Example 11 3.4 X ^{10 8} 18 0.27 40 630 769 20.1 42.5 251 Inventive Example 12 3.2 X ^{10 8} 16 0.31 40 625 775 20.4 43.2 223 Inventive Example 13 2.8X10 ⁸ 19 0.35 40 638 784 20.1 42.1 220 Conventional Steel 1 35 406 438 - Conventional Steel 2 35 405 441 - Conventional Steel 3 25 681 629 - 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 25 475 532 - Conventional Steel 8 50 504 601 - Conventional Steel 9 60 526 648 - 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 5, the hot rolled material produced by the present invention has a composite structure of bainite + ferrite, and the number of precipitates (Ti-based nitrides) was significantly finer and the number of precipitates was significantly increased than that of the conventional materials. I can see it well.

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 Example 1 22 35 55 76 19 374 -75 332 -66 290 -60 Inventive Example 2 21 34 54 78 18 385 -78 351 -69 304 -62 Inventive Example 3 22 36 55 82 19 362 -73 334 -65 297 -61 Comparative Example 1 56 83 187 34 45 125 -41 46 -34 24 -28 Comparative Example 2 68 95 192 23 44 105 -45 32 -32 19 -25 Inventive Example 4 23 38 64 75 20 354 -70 328 -65 287 -59 Inventive Example 5 25 42 56 79 18 365 -72 334 -66 299 -60 Comparative Example 3 52 78 135 44 28 181 -46 87 -32 56 -28 Comparative Example 4 47 72 127 52 30 151 -45 80 -34 62 -26 Inventive Example 6 24 38 54 74 18 385 -75 354 -69 306 -62 Inventive Example 7 23 37 53 78 18 363 -72 337 -67 294 -60 Inventive Example 8 26 34 54 83 19 365 -73 339 -67 293 -60 Inventive Example 9 26 35 54 76 18 363 -73 330 -64 287 -59 Inventive Example 10 25 36 55 78 18 383 -72 345 -65 298 -61 Inventive Example 11 23 37 64 77 18 356 -71 328 -64 285 -59 Inventive Example 12 24 38 65 81 17 354 -71 321 -64 276 -58 Inventive Example 13 23 36 64 80 19 358 -72 320 -65 378 -59 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 degrees Celsius) 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 ℃)

As shown in Table 6, the welded heat affected zone austenite grain size at the maximum heating temperature of 1400 ℃, while the range of 54-64㎛ in the case of the present invention, while about 180 in the conventional material (4-6) It can be seen that the present invention examples are very excellent in inhibiting the austenite grains of the weld heat affected zone because the present invention is shown to have a range of μm or more. In the case of the invention, -40 ℃ impact toughness of the weld heat affected zone, which gave high heat input welding heat cycle with cooling time of 800-500 ℃ respectively, showed excellent toughness value of about 300J or more, and the transition temperature was below -60 ℃. The value shows excellent impact toughness. On the other hand, in case of the conventional materials, the impulse toughness of 0 ° C. was very low as 60 to 132 J. Therefore, it can be seen that the steels according to the present invention can significantly improve the impact toughness and the transition temperature of the weld heat affected zone compared to the existing steels.

As described above, the present invention has a useful effect of providing a structural steel material capable of improving the impact toughness of the heat input welding heat affected zone to the steel of bainite + ferrite.

Claims (16)

By weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1%, N: 0.008-0.03%, B: 0.0003-0.01 %, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7 , Which satisfies 6.5≤ (Ti + 2Al + 4B + V) / N≤14, is composed of the remaining Fe and other unavoidable impurities, and has a microstructure of 30-80% of bainite and the remaining 20μm of ferrite High strength welded structural steel with excellent weld heat affected zone toughness. The welding heat effect according to claim 1, wherein the steel contains 0.01 to 0.2% and satisfies 0.3 ≦ V / N ≦ 9 and 7 ≦ (Ti + 2Al + 4B + V) / N ≦ 17. High strength welded structural steel with excellent toughness. The steel material according to claim 1 or 2, wherein the steel comprises Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%. High strength welded structural steel with excellent toughness of weld heat affected zone, characterized in that it contains one or more selected ones. The method of claim 1 or claim 2, wherein the steel is a high-strength weld structure excellent in weld heat affected toughness, characterized in that one or two selected from the group of Ca: 0.0005-0.005%, Rem: 0.005 to 0.05%. Steel. According to claim 3, wherein the steel is Ca: 0.0005-0.005%, Rem: 0.005 to 0.05% of the high strength welded structural steel having excellent weld heat affected zone toughness, characterized in that it contains one or two selected from the group. 3. The high strength welded structure according to claim 1 or 2, wherein TiN precipitates of 0.01-0.1 μm are distributed in the steel material at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less. Steel. By weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1%, N: 0.008-0.03%, B: 0.0003-0.01 %, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.005% or less, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7 , Austenitic recrystallization zone after satisfying 6.5≤ (Ti + 2Al + 4B + V) / N≤14 and heating steel slab composed of remaining Fe and other unavoidable impurities at a temperature of 1150-1250 ℃ for 60-180 minutes Hot-rolled at a rolling ratio of 40% or more, followed by cooling at a rate of 5-20 ° C./sec to bainite transformation end temperature ° C. ± 10 ° C. Way. 8. The welding heat effect according to claim 7, wherein the steel contains 0.01 to 0.2% and satisfies 0.3≤V / N≤9 and 7≤ (Ti + 2Al + 4B + V) / N≤17. Method of manufacturing high strength welded structural steel with excellent toughness. The method of claim 7 or 8, wherein the steel is Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0% A method for producing a high strength welded structural steel with excellent weld heat affected zone toughness, characterized in that it contains one or more selected ones. [9] The high strength welded structure according to claim 7 or 8, wherein the steel contains one or two selected from the group of Ca: 0.0005-0.005% and Rem: 0.005 to 0.05%. Method of manufacturing steels. 10. The method of claim 9, wherein the steel contains at least one selected from the group consisting of Ca: 0.0005-0.005% and Rem: 0.005-0.05%. . According to claim 7, The steel slab is deoxidized element having a greater deoxidizing power than Ti in molten steel prior to the Ti input, deoxidized dissolved oxygen to less than 30ppm, and then added within 10 minutes so that the Ti is 0.005 ~ 0.2% after desorption A method of manufacturing a high strength welded structural steel with excellent toughness in welded heat affected zone, characterized in that it is continuously cast after gas treatment. The method of claim 12, wherein the deoxidation is performed by adding Mn, Si, and Al in order. 14. The welding heat according to claim 12 or 13, wherein the continuous casting is coldly cooled at a specific amount of 0.3 to 0.35 l / kg in a secondary cooling zone while casting molten steel at a speed of 0.9 to 1.1 m / min. Method for manufacturing high strength welded structural steel with excellent impact toughness. The welding is applied to the steel (base metal) of claim 6 to form a prior austenite of 80 μm or less in the weld heat affected zone, followed by quenching to 70% of the ferrite of 20 μm or less in the microstructure of the weld heat affected zone. High-strength welded structure with excellent toughness of the weld heat affected zone composed of the above normal percentage. 16. The high strength welded structure of claim 15, wherein the ferrite has a polygonal ferrite and needle-like ferrite grains.