KR100482188B1 - Method for manufacturing high strength steel plate having superior toughness in weld heat-affected zone by recrystallization controlled rolling - Google Patents

Abstract

The present invention relates to a method for manufacturing structural steel used in welded structures such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes, etc. The object of the present invention is to provide fine grains of ferrite and bainite by recrystallization control rolling and control cooling. The present invention provides a method of manufacturing a welded structural steel that can improve the toughness of the weld heat affected zone by increasing the number of TiN precipitates having excellent high temperature stability to the microstructure while improving the strength of the base metal by forming a composite structure. The present invention for achieving the above object, in the weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.005-0.1%, N: It is composed of 0.008-0.030%, 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, the remaining Fe and other unavoidable impurities, and 1.2≤Ti Steel slab that satisfies /N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14 in the temperature range of 1000 ~ 1200 ℃ After heating for 3 hours, the heated slab is cooled to a temperature range of Ar ₃ to (Ar ₃ + 50 ° C.), hot to a reduction ratio of 40% or more, and a temperature of (Ar ₁ +30) to (Ar ₁ + 80 ° C.). Recrystallization comprising cooling to a range at a rate of 5-15 ° C./sec, hot rolling at a reduction ratio of 50% or more, and cooling at a rate of 5-20 ° C./sec to a temperature of bainite transformation end temperature ± 10 ° C. Technical aspects of the method for manufacturing a high strength welded structural steel by controlled rolling are described.

Description

Method for manufacturing high strength steel plate having superior toughness in weld heat-affected zone by recrystallization controlled rolling}

The present invention relates to a method for manufacturing structural steel used in welded structures, such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes. More specifically, recrystallization-controlled rolling and controlled cooling create a composite structure of fine grained ferrite and bainite, improving the strength of the base metal, and increasing the number of TiN precipitates having excellent high temperature stability in the microstructure, thereby improving the toughness of the weld heat affected zone. It relates to a method of manufacturing a welded structural steel that can improve the.

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 in the case of welding a steel plate having a plate thickness of 25 mm or more, the above-described high heat input welding method capable of 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. Therefore, it is advantageous to enable high heat input welding in consideration of welding productivity. In other words, increasing the amount of heat input in the welding will be able to widen the range of use. The range of high heat input currently 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 is most degraded.

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, there is known a technique for delaying grain growth of the weld heat affected zone during welding by appropriately distributing an oxide or Ti-based carbonitride, which is stable at high temperature, to steel materials. Such techniques include Japanese Patent Laid-Open No. 11-140582, Hei 10-298708 Hei 10-298706 Hei 9-194990 Hei 9-324238 Hei 8-60292, (S) 60-245768, (Square) 5-186848, (S) 58-31065, (S) 61-79745, Journal of the Japan Welding Society, Vol. 52, No. 2, 49 and Japanese Patent Laid-Open And 64-15320.

Japanese Patent Laid-Open No. 11-140582 is a representative technique using TiN precipitates, and has a toughness of about 200J at 0 ° C when 100 J / cm of heat input (maximum heating temperature of 1400 ° C) is applied. Is about 300J). In this prior art, Ti / N is substantially managed at 4-12, so that TiN precipitates of 0.05 µm or less are 5.8 × 10 ³ pieces / mm 2 to 8.1 × 10 ⁴ pieces / mm 2, and TiN precipitates of 0.03 to 0.2 μm are 3.9. The toughness of the welded portion is secured by being precipitated at × 10 ³ pieces / mm 2 to 6.2 × 10 ⁴ pieces / mm 2.

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), Since the difference is about 100J, there is a limit in securing the reliability of the steel structure according to superheat input welding of thickened steel. In addition, as a method for securing a desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or more and quenched and then reheated for hot rolling, thereby increasing the manufacturing cost due to two heat treatments. In addition, the base material has a poor ferrite + pearlite structure, yield strength of 400 MPa, tensile strength of 530 MPa, and elongation of 20% or less.

Until now, many techniques have been known to improve the toughness of the weld heat affected zone during high heat input welding. However, the case of remarkably improving the toughness of the weld heat affected zone during superheated heat welding, which is maintained for more than 1350 ° C for a long time while securing the mechanical properties of the base metal, It has not been announced yet. In particular, there are few technologies in which the toughness of the weld heat affected zone shows the same level as that of the base metal. Therefore, if the above problems can be solved, super heat input welding of the thickened steel is possible, so that it is possible to achieve high efficiency of the welding operation as well as to secure the high-rise of the steel structure and the reliability of the steel structure.

The present invention improves the strength of the steel (base material) by making the microstructure a composite structure of fine grains of ferrite and bainite, while uniformly distributing fine TiN precipitates in the microstructure to minimize the difference in toughness of the base material and the weld heat affected zone. It is to provide a method of manufacturing a welded structural steel that can be, and its purpose is to.

The present invention for achieving the above object, in the weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.005-0.1%, N: It is composed of 0.008-0.030%, 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, the remaining Fe and other unavoidable impurities, and 1.2≤Ti Steel slab satisfying /N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14 is heated in the temperature range of 1000 ~ 1200 ℃. Next, the heated slab is cooled to a temperature range of Ar ₃ to (Ar ₃ + 50 ° C.) at a rate of 5 to 15 ° C./sec, first hot rolled at a reduction ratio of 40% or more, and Ar ₁ to (Ar ₁ +50). Cooling at a rate of 5-15 ° C./sec to a temperature range of 50 ° C.), hot rolling at a reduction ratio of 50% or more, and cooling at a rate of 5-20 ° C./sec to a temperature of bainite transformation end temperature ± 10 ° C. It is configured to include.

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 the heat input welding is applied to the steel, and the austenite formed in the manufacturing process of the steel (hot rolling process). Use for convenience to distinguish from.

The present inventors studied a method of simultaneously securing the toughness of the weld heat affected zone while improving the strength and toughness of the steel (base material), and the base material strength when the microstructure of the base material is a composite structure of bainite and fine grains of ferrite It was confirmed that it is very effective for the improvement and that even if fine TiN is uniformly distributed in these microstructures, the austenite grain size of the weld heat affected zone is about 80 μm or less, so that the toughness of the weld heat affected zone is not a problem. .

Starting from this point of view, the inventors of the present invention have been able to derive the following method for 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 the fine TiN precipitates, the solubility product showing stability of the precipitates at high temperatures is reduced.

[2] The recrystallization-controlled rolling refines the ferrite grains of the base material to manage the size of the old austenite of the weld heat affected zone to about 80 μm or less. By refining the ferrite grains of the steel material, the strength and toughness of the base material are simultaneously improved. Also,

[3] By reducing the Ti / N ratio, BN and AlN precipitates are effectively precipitated, thereby increasing the fraction of ferrite in the weld heat affected zone, and inducing the ferrite into needles or polygons that are effective for toughness improvement.

[4] The strength of the base metal is improved through accelerated cooling in the rolling process. These [1] [2] [3] [4] are demonstrated in more detail.

[1] management of TiN precipitates

When heat input welding is applied to steel materials, the heat affected zone near the melting line is heated to about 1400 ° C or more, and the TiN precipitates deposited in the base materials are partially dissolved or coarsened to suppress the austenite grain growth. The effect disappears.

The inventors noticed that this phenomenon is caused by the diffusion of TiN precipitates distributed in the base metal by the dissolution of the dissolved Ti atoms by the heat of welding. As a result of examining the characteristics of TiN precipitates according to the ratio of Ti / N, the high nitrogen environment (Ti / 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.

According to the present invention, when the ratio of Ti and N (Ti / N) is heated to a 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, resulting in the fine TiN precipitate having a size of 0.01-0.1 μm. A result of distributing more than 1.0 × 10 ⁷ / mm ² at intervals of 0.5 μm or less can be obtained. Increasing the nitrogen content at the same Ti content causes all Ti atoms to be easily combined with nitrogen atoms, and the amount of solid solution Ti decreases in a high nitrogen environment. The solubility was lowered.

Furthermore, the present inventors are interested in the process of refining steel, which is important to solidify most of Ti without forming Ti as an oxide in molten steel in order to uniformly distribute a large amount of fine TiN precipitates. As a result, it was confirmed that Ti was not formed as an oxide and dissolved in steel when the amount of dissolved oxygen, which greatly affects the formation behavior of oxide, was controlled to 30 ppm or less.

According to the present invention, by managing the deoxidation conditions so that Ti is dissolved in the steel as much as possible, and when the ratio of Ti / N is controlled, fine TiN precipitates of 0.01-0.1 μm may be 1.0 × ^{10 7} / mm 2 or more at intervals of 0.5 μm or less. It can be distributed.

[2] recrystallization-controlled rolling (fine grains in ferrite grains)

According to the present invention, it is important to refine the size of the ferrite in the microstructure of ferrite + bainite with the management of precipitates in order to make the size of the old austenite to 80 μm while improving the strength and toughness of the base material. will be. In the present invention, the refinement of the ferrite is achieved by using recrystallization control rolling. Recrystallization-controlled rolling uses the precipitate (TiN) during high-temperature rolling to suppress grain growth, grain refining by static recrystallization during rolling, and ferrite grains due to precipitation of carbides during transformation (WC, Fe ₃ C, VC, etc.) Make it angry.

[3] microstructure of weld heat affected zone

The facts of the present invention revealed that the toughness of the weld heat affected zone has a significant influence not only on the size of the austenite grains but also on the amount and shape of the ferrite precipitated at the former austenite grain boundaries when the base material is heated above 1400 ° C. Will be. 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 , BN precipitates are used for this purpose.

[4] bainite tissue fraction control

The inventors of the present invention confirmed that by controlling the accelerated cooling rate after hot rolling, a fine ferrite and an appropriate bainite structure fraction can be obtained. At this time, it was confirmed that the properties of the weld heat affected zone is not related to the change in the microstructure of the base material.

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

[Composition of 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 0.005-0.1%.

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 at least 0.005%, but if the content exceeds 0.1%, Weedman Stetten is vulnerable to toughness during the cooling of the weld 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 precipitates and Ti oxides 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, the content of TiN and AlN precipitates and the spacing of precipitates, the distribution of precipitates, the frequency of complex precipitation with oxides, the high temperature stability of the precipitates themselves, etc. have a significant influence, so 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. It is preferable to limit the upper limit to 0.03% because of this.

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, which are 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%.

If 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 adversely affects the toughness of the base metal. It is not desirable because it is crazy.

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 of the base material increases, which is detrimental to the toughness of the weld heat affected zone. On the other hand, if the Ti / N ratio is higher than 2.5, coarse TiN is crystallized in the molten steel in the steelmaking process to obtain a uniform distribution of TiN, and excess Ti remaining without solidification as TiN remains in a solid solution to weld heat. 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 zone formed as described above, V is further added, and the content of this 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 precipitate 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 is effective for improving the toughness of the weld heat affected zone, so that the content is 0.1-1.5%. 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. Exceeding 1.0% results in degradation of the base metal and HAZ toughness.

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 if 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.

[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.

A feature of the present invention is that in such a deoxidation process, the dissolved oxygen is adjusted to an appropriate level or lower, and then Ti is added, so that Ti is mostly 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. It does not affect the effect of.

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 process of playing, and if it contains more than 0.2%, the effect is saturated and 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 present invention, the continuous casting speed is about 0.9 to 1.2 m / min. The reason is that when the casting speed is less than 0.9m / min, it is advantageous for cast surface cracks, but productivity is lowered. When it is faster than 1.2m / min, cast surface cracks are more likely to occur.

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 (recrystallization controlled rolling)

In the present invention, it is characterized by setting the rolling conditions that can secure the appropriate fraction of bainite through control cooling while refining the ferrite grains through recrystallization control rolling.

In the present invention, the slab is heated to a temperature range of 1000 ~ 1200 ℃. If the temperature is less than 1000 ° C., the heating time for depositing the fine TiN precipitate increases, which is not preferable. At the heating temperature of more than 1200 ° C., the size of the fine TiN precipitate increases, which is not preferable. The heated slab is cooled to a temperature range of Ar ₃ to (Ar ₃ + 50 ° C.) and first hot rolled at a reduction ratio of 40% or more. Below Ar ₃ , it becomes a ferrite + austenite two-phase structure, and coarse structure in which coarse ferrite remains in a later process is not preferable to secure a homogeneous microstructure, and in the case of Ar ₃ + 50 ° C, austenite particles are too coarse. This is because the ferrite refining effect in the post-process is less, which degrades the mechanical properties of the steel. If the reduction ratio is less than 40%, the recrystallized austenite grain size becomes large.

After rolling as described above, it is cooled at a rate of 5-15 ° C./sec to a temperature range of (Ar ₁ + 30 ° C.) to (Ar ₁ + 80 ° C.), and then secondary hot rolling is performed at a rolling reduction rate of 50% or more. do. At a temperature below (Ar ₁ + 30 ° C.), the ferrite transformation amount increases, and thus the amount of austenite to be transformed into bainite is insufficient, thereby making it difficult to secure high strength. In the case of (Ar ₁ + 80 ° C.), the transformed ferrite precipitates a large amount of coarse ferrite and remains in a coarse stretched form after rolling, which has a detrimental effect on the refinement of ferrite. When the cooling rate is less than 5 ° C / sec it is difficult to ensure the fine ferrite grains during the abnormal reverse rolling, which is the next step of the present invention because it causes the grain growth of the recrystallized fine ferrite. This is because it is difficult to secure the austenite recrystallization region at a cooling rate exceeding 20 ° C / sec, and thus it is difficult to secure fine ferrite grains during abnormal reverse rolling. In addition, when the amount of reduction is less than 50%, since the amount of strained organic ferrite transformation is small, coarse ferrite is formed, which is not preferable for securing fine ferrite.

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 as described above, control cooling is carried out in the range of 5-20 ° C / sec to bainite transformation end temperature (about 450 ° C) ± 10 ° C. Since the phase transformation of the present invention steel does not occur in a section other than the bainite transformation end temperature ± 10 ° C, the control cooling is limited to this section. If the controlled 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 if it exceeds 20 ° C / sec, the martensite phase ratio increases to be harmful to the base material toughness.

After cooling as above, it is cooled to room temperature. Of course, when the bainite transformation is completed, water cooling may be performed without air cooling to room temperature. However, the austenite remaining after transformation into ferrite and bainite may go to martensite.

[Organization of Steel]

· Microstructure

In the present invention, after the hot rolling, the steel microstructure is a composite structure of ferrite + bainite, and the structure fraction of bainite is preferably in the range of 30-80%. If less than 30% it is difficult to secure the appropriate base material strength for showing the effect of the present invention, in the case of more than 80% it is difficult to secure the base material toughness.

And it is preferable to make the size of the ferrite crystal grain of a base material into 8 micrometers or less. As the ferrite grains become finer, the toughness of the weld heat affected zone can be secured by atomizing the old austenite grains in the weld heat affected zone during the high heat input welding.

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 TiN to 0.01-0.1 μm and 1.0 × 10 ⁷ or more per 1 mm ² . The reason is that less than 0.01 μm, it is easily re-applied to the base metal during large heat input welding, and the effect of suppressing the growth of austenite grains is insufficient. If the thickness exceeds 0.1 μm, pinning of the austenite grains is achieved. Less effective) and behaves like coarse non-metallic inclusions, detrimentally affecting 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.

[Welding Structure]

The welded steel material provided according to the present invention is a composite structure of bainite + ferrite, and the grain size of the ferrite is 8 µm or less. In addition, TiN precipitates have a size of 0.01-0.1 μm and 1.0 × 10 ⁷ or more per 1 mm ² , and the TiN precipitates are finely distributed at intervals of 0.5 μm or less.

When the high heat input welding is applied to such steels, the old austenite grain size is 80 μm or less. If the austenite grain size is greater than 80 µm, low-temperature structure (martensite or upper bainite) is easily formed due to the increase in hardenability, which is detrimental to the weld heat affected zone toughness. In addition, even if ferrites having different nucleation sites are produced at the old austenite grain boundaries, ferrite coalesces during grain growth, which adversely affects toughness. Therefore, the austenite grain size of the weld heat affected zone for showing the effect of the present invention is 80 μm or less.

When the high heat input welding is applied and quenched as described above, the ferrite having the microstructure size of the heat-affected portion of about 20 μm or less has a phase fraction of about 70% or more. If 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. The ferrite of the present invention is more advantageous for toughness when it has the characteristics of about 40% or more of polygonal ferrite and acicular ferrite.

In the hot rolling process of the present invention, hot charge rolling and direct rolling, which are well known according to a user's use, may be applied, and various techniques such as well-known control rolling and control cooling may be applied. In addition, in order to improve the mechanical properties of the hot rolled sheet produced according to the present invention, heat treatment may be applied. Thus, even if the known techniques are applied to the present invention, it is natural to interpret that this is a simple change of the present invention substantially within the technical idea of the present invention.

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

EXAMPLE

Steel grades having the composition as shown in Table 1 were prepared as slabs according to the conditions of Tables 2 and 4 by melting in a converter, and the composition ratios between the alloying elements of the steel types according to the conditions of the present invention are shown in Table 3 to show the effects of the present invention. .

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

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 weld heat affected zone, it is heated to the maximum heating temperature (1200 ~ 1400 ℃) to 140 ℃ / sec condition by using a simulated welding simulator (1). Hold for a second and then quenched with He gas. The quenched specimens were ground and corroded to determine the austenite grain size at the highest heating temperature condition by KS (KS D 0205).

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, processed into an impact test piece, and evaluated through a Charpy impact test at -40 ° C.

Chemical composition (% by weight) C Si Mn P S Al Ti B (ppm) N (ppm) W Cu Ni Cr Mo Nb V Ca REM O (ppm) Invention 1 0.12 0.13 1.54 0.006 0.005 0.04 0.014 7 120 0.005 - - - - - 0.01 - - 11 Invention 2 0.07 0.12 1.50 0.006 0.005 0.07 0.05 10 280 0.002 - 0.2 - - - 0.01 - - 12 Invention 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015 3 110 0.003 0.1 - - - - 0.02 - - 10 Invention 4 0.10 0.12 1.48 0.006 0.005 0.02 0.02 5 80 0.001 - - - - - 0.05 - - 9 Invention 5 0.08 0.15 1.52 0.006 0.004 0.09 0.05 15 300 0.002 0.1 - 0.1 - - 0.05 - - 12 Invention 6 0.10 0.14 1.50 0.007 0.005 0.025 0.02 10 100 0.004 - - - 0.1 - 0.09 - - 9 Invention 7 0.13 0.14 1.48 0.007 0.005 0.04 0.015 8 115 0.15 0.1 - - - - 0.02 - - 11 Invention Material 8 0.11 0.15 1.52 0.007 0.005 0.06 0.018 10 120 0.001 - - - - 0.015 0.01 - - 10 Invention 9 0.13 0.21 1.50 0.007 0.005 0.025 0.02 4 90 0.002 - - 0.1 - - 0.02 0.001 - 12 Invention 10 0.07 0.16 1.45 0.008 0.006 0.045 0.025 6 100 0.05 - 0.3 - - 0.01 0.02 - 0.01 8 Inventive Steel 11 0.09 0.12 1.48 0.006 0.003 0.047 0.019 10 130 0.01 - 0.1 - - - - - - 12 Conventional Steel 1 0.05 0.13 1.31 0.002 0.006 0.0014 0.009 1.6 22 - - - - - - - - - 22 Conventional Steel 2 0.05 0.11 1.34 0.002 0.003 0.0036 0.012 0.5 48 - - - - - - - - - 32 Conventional Steel 3 0.13 0.24 1.44 0.012 0.003 0.0044 0.010 1.2 127 - 0.3 - - - 0.05 - - - 138 Conventional Steel 4 0.06 0.18 1.35 0.008 0.002 0.0027 0.013 8 32 - - - 0.14 0.15 - 0.028 - - 25 Conventional Steel 5 0.06 0.18 0.88 0.006 0.002 0.0021 0.013 5 20 - 0.75 0.58 0.24 0.14 0.015 0.037 - - 27 Conventional Steel 6 0.13 0.27 0.98 0.005 0.001 0.001 0.009 11 28 - 0.35 1.15 0.53 0.49 0.001 0.045 - - 25 Conventional Steel 7 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 - 0.3 - - - 0.036 - - - - Conventional Steel 8 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 - 0.32 0.35 - - 0.013 - - - - Conventional Steel 9 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 - - - 0.21 0.19 0.025 0.035 - - - Conventional Steel 10 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 - - 1.09 0.51 0.36 0.021 0.021 - - - Conventional Steel 11 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 - - - - - - 0.069 - - - Conventional steels (1, 2, 3) are invention steels (5, 32, 55) of Japanese Patent Application Laid-Open No. 9-194990, and conventional steels (4, 5, 6) are inventions of Japanese Patent Application Laid-Open No. 10-298708. Steels (14, 24, 28) and conventional steels (7, 8, 9, 10) are invention steels (48, 58, 60, 61) of JP-A-8-60292. Inventive Steel F of Patent Publication No. 11-140582

Steel grade used division Primary deoxidation sequence Dissolved oxygen after Al addition (ppm) Ti addition after finishing deoxidation (%) Casting speed (m / min) Specific quantity (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 19 0.014 1.1 0.35 Invention 2 Mn → Si 18 0.014 1.1 0.35 Invention 3 Mn → Si 18 0.014 1.1 0.35 Comparative Material 1 Mn → Si 25 0.014 0.5 0.15 Comparative Material 2 Mn → Si 21 0.014 1.8 0.75 Inventive Steel 2 Invention 4 Mn → Si 16 0.05 1.1 0.35 Invention Steel 3 Invention 5 Mn → Si 15 0.015 1.2 0.35 Inventive Steel 4 Invention 6 Mn → Si 15 0.02 1.2 0.35 Inventive Steel 5 Invention 7 Mn → Si 12 0.05 1.2 0.35 Inventive Steel 6 Invention Material 8 Mn → Si 17 0.02 1.2 0.35 Inventive Steel 7 Invention 9 Mn → Si 18 0.015 1.0 0.35 Inventive Steel 8 Invention 10 Mn → Si 14 0.018 1.0 0.35 Inventive Steel 9 Invention 11 Mn → Si 19 0.02 1.1 0.35 Inventive Steel 10 Invention Material12 Mn → Si 23 0.025 1.1 0.35 Inventive Steel 11 Invention Material 13 Mn → Si 21 0.019 1.1 0.35 The manufacturing conditions of the conventional steel (1-11) are not specifically described.

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 Invention 1 1.2 17.1 3.3 0.8 8.9 Invention 2 1.2 17.1 3.3 0.8 8.9 Invention 3 1.2 17.1 3.3 0.8 8.9 Comparative Material 1 1.2 17.1 3.3 0.8 8.9 Comparative Material 2 1.2 17.1 3.3 0.8 8.9 Invention 4 1.8 28.0 2.5 0.4 7.3 Invention 5 1.4 36.7 5.5 1.8 14.2 Invention 6 2.5 16.0 2.5 6.3 14.0 Invention 7 1.7 20.0 3.0 1.7 9.5 Invention Material 8 2.0 10.0 2.5 9.0 16.4 Invention 9 1.3 14.4 3.5 1.7 10.3 Invention 10 1.5 12.0 5.0 0.8 12.7 Invention 11  2.2 22.5 2.8 2.2 10.2 Invention Material12 2.5 16.7 4.5 2.0 13.7 Invention Material 13 1.5 13.0 3.6 - 9.0 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 Heating temperature (℃) Heating time (min) 1st rolling Cooling rate (℃ / sec) Secondary rolling Cooling rate up to 450 ℃ ± 10 ℃ (℃ / sec) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount (%) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount (%) Invention 2 Inventive Example 1 1100 120 940 920 50 5 780 760 60 15 Inventive Example 2 1100 120 940 920 50 5 780 760 60 15 Inventive Example 3 1100 120 940 920 50 5 780 760 60 15 Comparative Example 1 1100 120 1050 1030 40 5 780 760 50 2 Comparative Example 2 1100 120 940 920 30 5 780 760 40 45 Invention 1 Inventive Example 4 1050 130 950 920 45 8 790 770 55 15 Invention 3 Inventive Example 5 1050 130 950 920 45 6 790 770 55 15 Comparative Material 1 Comparative Example 3 950 50 940 930 30 8 880 860 50 16 Comparative Material 2 Comparative Example 4 1300 250 1050 1000 40 7 900 880 50 15 Invention 4 Inventive Example 6 1050 150 950 920 45 7 780 765 55 16 Invention 5 Inventive Example 7 1050 150 940 920 50 7 780 765 65 15 Invention 6 Inventive Example 8 1100 120 940 920 50 7 780 760 55 15 Invention 7 Inventive Example 9 1100 130 960 930 50 8 780 760 55 15 Invention Material 8 Inventive Example 10 1100 130 950 920 50 8 790 770 55 15 Invention 9 Inventive Example 11 1150 130 960 930 50 8 790 770 60 15 Invention 10 Inventive Example 12 1150 130 960 930 50 8 790 765 60 15 Invention 11 Inventive Example 13 1050 150 950 920 45 8 790 770 60 15 Invention Material12 Inventive Example 14 1100 130 940 920 45 8 790 770 60 16 Invention Material 13 Inventive Example 15 1150 120 940 920 50 8 790 770 60 16 Conventional Steel 11 1200 - Ar ₃ or higher 80 Cooling Manufacturing conditions of conventional steel (1-10) are not specifically presented

As shown in Table 5, the number of precipitates (Ti-based nitride) of the hot rolled material produced by the present invention has a range of 1.0 × ^{10 8} / mm ² or more, whereas the precipitates of conventional materials are 4.7X10 ⁵ / mm ^Since showing the range of ² or less, it can be seen that the number of the invention material is significantly increased while the invention material has a very fine precipitate size compared to the conventional material. In addition, it can be seen that the ferrite grain size (FGS) of the present invention steel is very fine in the structure of the base metal structure of about 8 μm or less, and all of the bainite phase fraction is 30% or more.

As shown in Table 6, the austenitic grain size at the maximum heating temperature of 1400 ° C., such as the weld heat affected zone, ranges from 52 to 65 μm in the case of the present invention. You can see that we go to the range. Therefore, it can be seen that the present invention steel is very excellent in inhibiting the austenite grains of the weld heat affected zone during welding. In addition, the ferrite phase fraction of the steel of the present invention was composed of about 70% or more at a welding heat input amount of 100 kJ / cm.

On the other hand, when comparing the impact toughness of the weld heat affected zone during high heat input welding, in the case of the present invention under the high heat input weld heat input condition of 250 kJ / cm (180-500 ° C. cooling time of 180 seconds), The impact toughness showed excellent toughness value of about 280J or more, and the transition temperature also showed a value of about -60 ° C. or less, indicating that the high heat input welding heat affected zone impact toughness was excellent. On the other hand, in the case of conventional steel, the impact toughness of 0 ° C. was very low as 60-132J. Therefore, it can be seen that the steels according to the present invention can significantly improve the impact toughness and transition temperature of the welded heat affected zone compared to the existing steels.

As described above, the present invention can further improve the impact toughness of the welded heat affected zone by further distributing a large amount of TiN precipitates stable at high temperature, and managing the microstructure of the base material as bainite + ferrite. As a result, the substrate strength is also useful to provide improved welded structural steel.

Claims (7)

By weight C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.005-0.1%, N: 0.008-0.030%, 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, composed of the remaining Fe and other unavoidable impurities, 1.2≤Ti / N≤2.5, 10≤N / B Steel slabs satisfying ≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14 were heated to a temperature range of 1000 to 1200 ° C, and the heated slabs were then heated to Ar _3- ( Cool down to a temperature range of Ar ₃ + 50 ° C.) and hot rolling at a reduction ratio of 40% or more, and at a rate of 5-15 ° C./sec to a temperature range of (Ar ₁ + 30 ° C.) to (Ar ₁ + 80 ° C.). Method for producing a high strength welded structural steel by recrystallization controlled rolling comprising cooling and hot rolling at a reduction ratio of 50% or more, and cooling at a rate of 5-20 ° C / sec to a temperature of bainite transformation end temperature ± 10 ° C. . The method of claim 1, wherein the steel material contains 0.01 to 0.2%, and the ratio of V and N (V / N) is 0.3 to 9, and 7≤ (Ti + 2Al + 4B) / N≤17. A method for producing a high strength welded structural steel by recrystallization-controlled rolling. 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%. A method for producing a high strength welded structural steel by recrystallization-controlled rolling, characterized by containing one or two or more selected from one or more selected from the group of Ca: 0.0005-0.005% and REM: 0.005 to 0.05%. The method of claim 1, wherein the steel slab is added to the molten steel prior to Ti injecting a deoxidation element having a greater deoxidizing power than Ti to deoxidize dissolved oxygen of the molten steel to less than 30ppm, and then added within 10 minutes so that the Ti is 0.005 ~ 0.2% A method for manufacturing a high strength welded structural steel by recrystallization-controlled rolling, characterized in that it is made by continuous casting of degassed molten steel. The method for manufacturing a high strength welded structural steel sheet according to claim 1 or 4, wherein the deoxidation is performed in the order of Mn, Si, and Al. The continuous casting of the steel slab according to claim 1 or 4, wherein the continuous casting of molten steel at a rate of 0.9 to 1.2 m / min, while the secondary slab is weakly cooled at a specific amount of 0.3 to 0.35 l / kg in a secondary cooling zone. A method for producing a high strength welded structural steel by recrystallization-controlled rolling. The steel material according to any one of claims 1, 2, and 4, wherein the steel has a microstructure of 30-80% of bainite and a remainder of 8 µm or less of ferrite, and a TiN of 0.01-0.1 µm. A method for manufacturing a welded structural steel with excellent toughness in weld heat affected zones in which precipitates are distributed at 1.0x10 ⁷ holes / mm 2 or more at intervals of 0.5 μm or less.