KR100368244B1 - Method for steel plate having superior toughness in weld heat-affected zone - 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 present invention is to use TiO oxide while securing fine TiN precipitates by sedimentation on low-nitrogen steel slabs. By providing a method for manufacturing a welded structural steel that can improve the toughness of the weld heat affected zone.

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: 0.005% or less, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, and other Fe and other impurities, and 4≤Ti / Making a low nitrogen steel slab satisfying O ≦ 10;

Heating the slab at a temperature of 1100 to 1250 ° C. for 60 to 180 minutes to immerse the steel to satisfy the following relationship with Ti, B, and Al in the range of 0.008 to 0.03%; And

1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14

The heated slab is hot rolled in an austenite recrystallization zone at a rolling ratio of 40% or more, and then cooled to a ferrite transformation end temperature ± 10 ° C at a rate of 1 ° C / min or more. The technical subject matter is a method for manufacturing structural steels.

Description

Method for manufacturing welded structural steel with excellent toughness in welding heat affected area {Method for steel plate having superior toughness in weld heat-affected zone}

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 present invention relates to a method for producing a welded structural steel material which can improve the toughness of the weld heat affected zone by using TiO oxide while uniformly distributing a large amount of fine TiN precipitates by immersion in a low nitrogen steel slab. .

Recently, the steel used in accordance with the trend of high-rise building, structure has been replaced by thick steel. In order to weld such thick steels, high-efficiency welding is inevitable. As a technique for welding thickened steels, a high-pass heat submerged welding method and an electro-welding method capable of one-pass welding are widely used. In addition, in the field of shipbuilding and bridges, even 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. Therefore, it is advantageous to enable high heat input welding in consideration of welding productivity. In other words, increasing the amount of heat input in the welding will be able to widen the range of use. The range of high heat input currently in use is about 100-200kJ / cm, but in order to weld steel materials with a thicker plate thickness of 50mm or more, it is possible to have a super heat input range of 200-500kJ / 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 is the site where the toughness of the weld deteriorates most.

Therefore, in order to secure the stability of the welded structure, it is necessary to suppress the 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 disclosed 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 a high temperature, to steel materials. For example, Japanese Patent Application 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 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 making the ferrite fine by depositing 10 ³ pieces / mm 2 to 6.2 × 10 ⁴ pieces / mm 2.

However, according to this prior art, when the 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), and also the base material. Since the toughness difference between the and the heat affected zone is about 100J, there is a limit in securing the reliability of the steel structure due to superheated welding of the thickened steel. In addition, as a method for securing the desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or higher and quenched, followed by a reheating process for hot rolling. Becomes Moreover, in this prior art, since the high nitrogen molten steel containing 0.005-0.02% of N is continuously cast into an ingot, the surface crack of a cast steel is high. That is, since N is an austenite stabilizing element and austenite is maintained for a long time in the solidification process of the ingot, impurity elements such as P and S cause segregation in the uncoagulated portion to cause cast cracks.

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

Until now, many techniques have been known to improve the toughness of the weld heat affected zone during high heat input welding using TiN precipitates and oxides such as Al-based or MgO. However, the toughness of the weld heat affected zone during superheated heat welding maintained at 1350 ° C. or more for a long time is known. No significant improvement has yet been announced. In particular, there are few technologies in which the toughness of the weld heat affected zone shows an equivalent level to 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.

According to the present invention, while the low nitrogen steel slab is immersed in the slab surface of the high nitrogen steel, the difference between the toughness of the base material and the heat affected zone is minimized even in the heat input range from the medium heat input to the super heat input, while fundamentally blocking the occurrence of cracks. To provide a method for producing a welded structural steel with excellent toughness of the heat affected zone.

Method for producing a welded structural steel 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: 0.005% or less, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, with the remaining Fe and other impurities Making a low nitrogen steel slab which is formulated and satisfies 4 ≦ Ti / O ≦ 10;

Heating the slab at a temperature of 1100 to 1250 ° C. for 60 to 180 minutes to immerse the steel to satisfy the following relationship with Ti, B, and Al in the range of 0.008 to 0.03%; And

1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14

The heated slab is hot rolled in an austenite recrystallization zone with a rolling ratio of 40% or more, and then cooled to a ferrite transformation end temperature ± 10 ° C at a rate of 1 ° C / min 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 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 for improving the toughness of the weld heat affected zone while preventing surface cracks generated in high nitrogen steel. As a result, instead of making steel slabs from molten nitrogen steel, it has excellent high temperature stability through nitriding in a subsequent process. It was confirmed that the toughness of the weld heat affected zone would not be a problem if the TiN precipitates were uniformly distributed and the grain size of the old austenite was managed below the threshold (about 80 μm).

Based on these studies, in the present invention,

[1] with the use of TiN precipitates and TiO oxides, which are distributed by nitriding in low nitrogen steels,

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

[3] Effectively use TiO oxide, BN and AlN precipitates to increase the ferrite fraction in the weld heat affected zone, and enhance the toughness improvement effect by promoting the transformation of polygonal or needle-like ferrite in Guostenite. These [1] [2] [3] are demonstrated in more detail.

[1] TiN Precipitate and TiO Oxide Management by Nitriding

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 the TiN precipitates are improved. Even more interesting is the ratio of Ti and N (Ti / N), even though steel slabs are made of low nitrogen steel with less than 0.005% low probability of cast surface cracking and made into high nitrogen steel by immersion treatment in slab heating furnace during the rolling process. In the range of 1.2 to 2.5, the amount of solid solution Ti is extremely reduced and the high temperature stability of TiN precipitates is increased, so that the fine TiN precipitates of 0.01-0.1㎛ size are 1.0x10 ⁷ / mm2 or more at intervals of 0.5㎛ or less. Distribution results were obtained. Increasing the nitrogen content at the same Ti content, all the Ti atoms dissolved are easily combined with the nitrogen atom, and in the high nitrogen environment, the amount of solid solution Ti decreases, which makes TiN precipitates more stable at higher temperatures than in the case of low nitrogen content. It is analyzed that this is because the solubility product is lowered. In view of the fact that the present invention can promote 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.

In addition, in the present invention, TiO is managed at 4-10 to precipitate and distribute TiO oxides that are stable even at 1400 ° C. or higher, thereby suppressing the growth of old austenite grains.

[2] ferrite grain size management

According to the study of the present invention, it is important that the size of the ferrite is 20 μm or less while the microstructure of the base material is a composite structure of ferrite + pearlite in order to make the size of the old austenite at 80 μm in the weld heat affected zone. At this time, the refinement of the ferrite is to inhibit the growth of the ferrite grains generated in the cooling process after hot rolling using carbides (VC, WC) as well as the austenite grain refinement by the steel processing during hot rolling.

[3] microstructure of weld heat affected zone

In the present invention, the Ti / O ratio is controlled to 4-10 to appropriately proportion the ratio of Ti contained in the oxide, and a large amount of fine ferrite is produced in the former austenite mouth using precipitates such as AlN and BN. Most ferrites are polygonal ferrites and needle-like ferrites, which greatly improve the toughness of heat-affected zones.

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

[Welding Structural Steels]

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

When the content of carbon (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, are transformed during cooling to lower the low temperature impact toughness of structural steel, and also the hardness of the welded portion. Or increasing the strength resulting in deterioration of toughness and generation of weld cracks.

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

If the content of silicon is less than 0.01%, the deoxidation effect of molten steel is insufficient during steelmaking and the corrosion resistance of steel is reduced. If the content is more than 0.5%, the effect is saturated, 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 that improves deoxidation, weldability, hot workability and strength in steel, and precipitates in the form of MnS around Ti-based oxides, which affects the formation of acicular and polygonal ferrites effective for improving the toughness of weld heat affected zones. Element. Such Mn forms a solid solution to form a solid solution in the matrix to strengthen the matrix to secure the strength and toughness, for this purpose it is preferred to add more than 0.4%. However, when the Mn content exceeds 2.0%, tissue heterogeneity due to Mn segregation has a detrimental effect on the toughness of the weld heat affected zone, rather than the effect of strengthening the solid solution. In addition, when Mn content is added more than 2.0%, macro segregation and micro segregation occurs depending on the segregation mechanism during steel coagulation, which promotes the formation of central segregation zone in the center of rolling, thereby creating low temperature transformation structure in the center of the base metal. do.

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

Al is an element necessary as a deoxidizer and forms an Al oxide by reacting with oxygen to prevent Ti from reacting with oxygen, which not only helps Ti to form fine TiN precipitates, but also is effective for forming fine AlN precipitates in steel. . In order to form fine AlN precipitates, the Al content is preferably made 0.0005% or more. However, if it exceeds 0.1%, AlN is precipitated and the remaining solid Al promotes the formation of Widmanstatten ferrite and phase martensite, which are vulnerable to toughness during the cooling process of the weld heat affected zone. Lowers.

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. It is preferable to add more than 0.005% of Ti in order to obtain such a fine TiN precipitation effect, but when it exceeds 0.2%, coarse TiN precipitates and Ti oxides are formed in molten steel, and thus it is impossible to suppress the growth of the austenite grains of the weld heat affected zone. Because it is not desirable.

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

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

The content of 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 is possible to minimize the growth of the old austenite grains in the weld heat affected zone during the high heat input welding and to increase TiN, AlN, BN, VN, NbN, etc. Increase the amount of precipitate; In particular, the size and precipitate spacing of the TiN and AlN precipitates, the distribution of precipitates, the frequency of complex precipitation with oxides, the high temperature stability of the precipitate itself, etc. have a significant effect, for this purpose it is preferably contained more than 0.008%. Therefore, it is preferable to make the slab from low-nitrogen molten steel of 0.005% or less, which is less likely to cause cracks on the surface of cast steel, and to make N content 0.008% or more through nitriding in the subsequent process. However, when the nitrogen content exceeds 0.03%, the effect is saturated, and toughness decreases due to an increase in the amount of solid solution nitrogen distributed in the weld heat affected zone, and it may be incorporated into the weld metal due to dilution during welding, resulting in a decrease in the toughness of the weld metal. Can be.

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

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

The content of 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. In particular, sulfur is precipitated in the form of MnS around Ti-based oxides, which affects the formation of needle-shaped and polygonal ferrites, which are effective for improving the toughness of welding heat affected zones. It is desirable to limit it to 0.03% or less.

The content of oxygen (O) is preferably limited to 0.0020-0.03%.

O is indispensable for forming Ti oxide by reacting with Ti in molten steel. Ti oxide promotes the transformation of acicular ferrite during ferrite transformation from guustenite in the weld heat affected zone. If the O content is less than 0.0020%, it is not preferable because the number of Ti oxides is small, and if it exceeds 0.03%, coarse Ti oxides and other oxides such as FeO are produced, which is not preferable.

The steel of the present invention has a sedimentation mass of nitrogen of 1.2-2.5 of Ti / N, 10-40 of N / B, 2.5-7 of Al / N, and 6.5- of (Ti + 2Al + 4B) / N. It is preferable to adjust to 14.

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 increase the amount of TiN, that is, the number of TiN precipitates. In other words, if the nitrogen content is increased at the same Ti content, all the Ti atoms dissolved in the cooling process during the playing process combine with the nitrogen atom, thereby increasing the fine TiN precipitation. Second, TiN is stable at high temperatures. That is, since the solubility product which shows the stability of the precipitate at high temperature such as the weld heat affected zone becomes smaller, the precipitate such as high nitrogen TiN is more stable than the case where the nitrogen content is low. On the other hand, 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. Also, excess Ti remaining without precipitation as TiN remains in a solid solution to weld heat. It is not preferable because it adversely affects the impact toughness. In addition, when the Ti / N ratio is less than 1.2, the amount of solid solution nitrogen in the base metal increases, which is detrimental to the toughness of the welding heat-facing portion.

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 that promotes the ferrite transformation of polygons at the old austenite grain boundary during the post-weld cooling process is insufficient, and when the N / B ratio is over 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 and the polygonal shape in the grain boundary Insufficient size and number of distribution of TiN, AlN, BN, and VN precipitates for ferrite formation and control of tissue fraction, and the effect is saturated when (Ti + 2Al + 4B) / N is more than 14 seconds. If V is added, the ratio of (Ti + 2Al + 4B + V) / N is preferably 7-17.

The ratio of Ti / O is preferably set to 4-10.

If the Ti / O ratio is less than 4, the number of TiO oxides required to inhibit austenite grain growth is insufficient, and the Ti ratio in TiO oxides becomes small, thus losing the function of acicular ferrite nucleation sites and improving the toughness of the weld heat affected zone. The effective acicular ferrite phase fraction is lowered. If the ratio of Ti / O is more than 10, the effect of inhibiting the weld heat affected zone austenite grain size crystallization is saturated, and the ratio of components such as Mn contained in the oxide becomes rather small, thus losing the function of the nucleation site of the ferrite in the mouth.

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.

In addition, it is preferable that the ratio of V / N after a immersion process shall be 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 on TiN + MnS precipitates for improving the toughness of the weld heat affected zone. If the V / N ratio exceeds 9, the size of the VN precipitates deposited at the TiN + MnS precipitate boundary is coarsened, and thus the VN precipitation frequency deposited at the TiN + MnS complex precipitate boundary is reduced, which is effective for the toughness of the weld heat affected zone. Reduce ferrite phase percentage.

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

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

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

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

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

Nb is an effective element from the viewpoint of securing the strength of the base material. For this purpose, Nb is added in an amount of 0.01% or more. However, Nb is undesirable because it causes coarse precipitation of coarse NbC, which is detrimental to the toughness of the base material.

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

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

Molybdenum (Mo) is preferably 0.05-1.0%.

Mo is also an element that increases the hardenability and improves the strength. The content thereof is 0.05% or more for securing the strength, but the upper limit is set to 1.0% like Cr for suppressing the HAZ hardening and the welding low temperature crack.

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

Ca and REM form an oxide having excellent high temperature stability, thereby suppressing the growth of the austenite grains 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.

[Manufacturing method]

Refining Process (Deoxidation and Degassing Process)

In general, the steel refining process is first refined in the refining furnace (electric furnace, electric furnace), and then the molten steel of the refining furnace is pulled out by the ladle (LF) to the secondary refining (external refining). After out-of-furnace refining, degassing (RH process) is performed. Normal deoxidation takes place in primary and secondary refining.

In the present invention, the dissolved oxygen is adjusted to an appropriate level in such a deoxidation process and then Ti is added to obtain a distribution of Ti oxide at an appropriate level. Based on the fact that the amount of dissolved oxygen greatly affects the production behavior of the oxide, the present inventors have found that there is an adequate amount of dissolved oxygen for producing a large number of oxides as a result of obtaining a desired Ti oxide. The value was confirmed to be 50-200ppm. In order to obtain this level of dissolved oxygen, a deoxidizer, which has a stronger deoxidizing power than Ti, must be added prior to the introduction of Ti to prevent the formation of coarse Ti oxide in molten steel while increasing the number of Ti oxides during the play. These Ti oxides can increase the precipitation frequency of precipitates, such as TiN, MnS, CuS, VN.

The deoxidizing power of the deoxidizer is as follows.

Cr <Mn <Si <Ti <Al <Rem <Zr <Ca ≒ Mg

The deoxidizer with greater affinity with oxygen has a higher rate of binding to oxygen in molten steel and a stronger binding force. Therefore, before adding Ti, a deoxidizer with a greater deoxidizing power than Ti must be added to ensure proper dissolved oxygen (50-200 ppm) to shorten the progress of the deoxidation process, and to secure proper dissolved oxygen to the desired TiO oxide. While forming Ti can be made into a TiN precipitate. That is, since the deoxidizer having a higher deoxidizing power than Ti is pre-injected, the ratio of Ti remaining in solid solution becomes considerably higher than that of coarse primary oxides in molten steel, so that a fine amount of Ti oxide and TiN nitride during solidification Will be able to secure.

In the present invention, before the addition of the elements (Al, REM Zr, Ca, Mg) having a greater deoxidizing power than Ti, by deoxidizing by adding elements such as Mn, Si, and then adding a deoxidizer having a higher deoxidizing power than Ti such as Al It is also preferable. This method has the advantage of easy control of the appropriate dissolved oxygen amount. As an example of the deoxidation operation pattern in the present invention, Mn, Si, Al in the first refining and then injecting Al in the second refining and Ti is added, this addition of Ti is vacuum degassing in the second refining In the case of adopting a gas treatment step, it is also possible to inject it after vacuum degassing.

In the present invention, it is preferable to add a deoxidation element having a greater deoxidizing power than Ti before deoxidizing the dissolved oxygen amount to 50-200 ppm. The reason is that when the dissolved oxygen amount is less than 50 ppm, the amount of oxygen in the molten steel is too small to form an adequate amount of Ti-based oxide for showing the effect of the present invention. In addition, it is difficult to obtain an alloying system composed of accurate chemical components by oxidizing other alloying elements.

Subsequently, it is preferable to add Ti so that the content is 0.005-0.2%, which improves the toughness of the weld heat affected zone by promoting the acicular ferrite transformation as well as the proper distribution of Ti oxide and Ti precipitates for showing the effect of the present invention. It is the range which considered the effective precipitate behavior of MnS, CuS, BN, VN, etc. together.

As described above, after the addition of Ti, the molten steel is refined, at which time the refining is performed for 0.5-20 minutes. When the retention time of the molten steel is less than 0.5 minutes, the coarse Ti-based primary oxide does not have time to separate the flotation and remains after coagulation, and the average size thereof increases, and when it exceeds 20 minutes, the oxides are bonded to each other. Therefore, it is difficult to secure an appropriate amount of solid solution Ti in the molten steel to coarsen and exhibit the effects of the present invention. In the present invention, the addition of Ti is possible either before or after vacuum degassing.

Casting process

In the present invention, since molten steel is low nitrogen steel, the casting speed may be either high speed or low speed during continuous casting. In order to obtain favorable internal quality, it is preferable to make it into the range of 0.9-1.2 m / min.

Slab reheating process

In the present invention, by controlling the ratio of Ti and N in the steel slab through the nitriding treatment in the slab heating furnace, by increasing the amount of very fine TiN precipitates and reducing the amount of solid solution Ti in the weld heat affected zone during welding, Osbal A method of maximally suppressing Ostwald ripening was derived. In addition to the fundamentally avoiding the problem of slab surface cracks commonly found in high-nitrogen steels, the nitridation effect in slab heating furnaces is two more. The first is to increase the amount of fine TiN precipitates, and the second is to stabilize the fine precipitated TiN at high temperature. That is, by increasing the nitrogen content in the base material at the same Ti content through the nitriding treatment, all Ti atoms can be combined with the nitrogen atoms during the heat treatment in the slab heating furnace to increase the amount of fine TiN precipitates.

In the present invention, it is preferable to make the slab nitrogen concentration by 0.008-0.03% by immersing the slab while heating the slab at 1100-1250 ° C. for 60-180 minutes. Nitrogen should be contained more than 0.008% to secure an appropriate level of TiN deposition in the slab, but if it exceeds 0.03%, the amount of nitrogen deposited on the surface of the slab increases more than the amount of nitrogen diffused into the slab and precipitated as fine TiN. As a result, hardening occurs on the surface of the slab, which affects the subsequent rolling process. At this time, if the slab heating temperature is less than 1100 ℃, the driving force for diffusion of the precipitated nitrogen is small, the number of fine TiN precipitates is small, and the heating time must be increased in order to increase the number of TiN precipitates, which increases the manufacturing cost cost There is. On the other hand, when the heating temperature is higher than 1250 ° C., the austenite grains of the slab grow during heating and affect the recrystallization during the rolling process. On the other hand, when the slab heating time is less than 60 minutes, the sedimentation effect is not exerted, and when the heating time is longer than 180 minutes, not only the cost of the unworking industry increases but also the growth of austenite grains in the slab affects the subsequent rolling process. It is not desirable because it is crazy.

When subjected to the nitriding treatment according to the present invention, the ratio of Ti / N in the slab is 1.2 to 2.5, the ratio of N / B is 10 to 40, the ratio of Al / N is 2.5 to 7, and the ratio of V / N is 0.3. It is preferable to impregnate N so that the ratio of -9 and (Ti + 2Al + 4B + V) N may be 7-17.

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.

After rolling as described above, in order to prevent ferrite grain growth, it is cooled at a rate of 1 ° C / min or more up to a ferrite transformation end temperature of ± 10 ° C. Preferably, it is cooled at a rate of 1 ° C./min or more to the ferrite transformation end temperature and thereafter air-cooled. Of course, it is possible to cool the ferrite transformation temperature below the end temperature, that is, above 1 ° C / min to room temperature, but it is uneconomical. If the cooling rate is less than 1 ° C / min, it causes a grain growth of the recrystallized fine ferrite, it is difficult to secure the ferrite grain size of the steel slab less than 20㎛.

· Microstructure of steel

In the present invention, after the hot rolling, the steel microstructure is made of ferrite + pearlite, and the size of the ferrite grains is preferably 20 μm or less. The reason is that when the grain size of the ferrite is larger than 20 μm, the size of the old austenite grains in 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.

In addition, the higher the percentage of ferrite in the composite structure of ferrite + perlite, the higher the toughness and elongation of the base metal, and the ferrite is 20% or more and most preferably 70% or more.

Distribution of precipitates

The former austenite grains of the weld heat affected zone are greatly influenced by the size, number and distribution of oxides or nitrides distributed in the base material when the austenite grain size of the base material is constant. In addition, since 30-40% of the nitrides distributed in the base material are welded to the base material at the time of high heat input welding (above the heating temperature of 1400 ℃ or more), the effect of inhibiting the growth of the austenite grains in the weld heat affected zone is reduced. There is a need for a uniform distribution of further nitrides taking into account the re-used nitrides. In order to suppress the growth of the old 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 for this is that less than 0.01 μm is easily re-used to the base metal during high heat input welding, and the effect of inhibiting the growth of the old austenite grains is insufficient. If the thickness exceeds 0.1 μm, pinning of the old 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 old austenite grains of the weld heat affected zone at the time of welding higher than the heat input to be 80 μm or less, which is a threshold value.

Oxide distribution

According to the present invention, it is preferable to control the ratio of Ti / O to 4-10 so that the particle size and the critical number of the Ti oxide in the base material are in the range of 0.5-2.0 μm and 1.0 × ^{10 2} -1.0x10 ³ per mm ² . The reason for this is that when the oxide size is less than 0.5 µm, the needle ferrite nucleation promoting effect is insufficient in the weld heat-affected grains. This is because the amount of transformation of the acicular ferrite is reduced. If the number of precipitates is 1.0X10 ² or less per 1 mm ^2, it is not preferable because the amount of acicular ferrite formed from Ti oxide is insufficient. If the amount of precipitate is more than 1.0X10 ^3, acicular ferrite formed from Ti oxide is formed from adjacent Ti oxide. It is not preferable because it overlaps with acicular ferrite and interferes with formation of acicular ferrite.

In the steel manufacturing method of the present invention, the slab can be produced by continuous casting or die casting. At this time, if the cooling rate is fast, because it is advantageous to finely disperse the precipitates, continuous casting with a high cooling rate is preferable, and the slab is advantageously thin in thickness for the same reason. After the slab is immersed in accordance with the present invention, in the hot rolling process, hot charge rolling and direct rolling, which are widely known according to the user's use, may be applied later. Various techniques such as cooling can 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 by a continuous casting method using samples as steel samples, and then hot rolled materials were prepared by immersion treatment and hot rolling under the conditions of Table 2. The composition ratio between the alloying element elements of this hot rolling material is shown in Table 3.

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 specimen 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). After holding, it was quenched with He gas. The quenched specimens were ground and corroded to determine the austenite grain size at the highest heating temperature condition by KS (KS D 0205).

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

Impact toughness evaluation of the welding heat affected zone is 800-500 ℃ cooling after heating the welding conditions corresponding to about 80 kJ / cm, 150 kJ / cm, 250 kJ / cm, that is, the maximum heating temperature to 1400 ℃ After the welding heat cycles of 60 seconds, 120 seconds, and 180 seconds were 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) Inventive Steel 1 0.11 0.23 1.54 0.006 0.005 0.04 0.014 7 45 0.005 - - - - - 0.01 - - 34 Inventive Steel 2 0.08 0.12 1.50 0.006 0.005 0.07 0.05 10 47 0.002 - 0.2 - - - 0.01 - - 54 Invention Steel 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015 3 42 0.003 0.1 - - - - 0.02 - - 32 Inventive Steel 4 0.11 0.13 1.49 0.006 0.005 0.02 0.02 5 40 0.001 - - - - - 0.05 - - 31 Inventive Steel 5 0.08 0.15 1.52 0.006 0.004 0.09 0.05 15 48 0.002 0.1 - 0.1 - - 0.05 0.001 - 52 Inventive Steel 6 0.10 0.14 1.50 0.007 0.005 0.025 0.02 10 38 0.004 - - - 0.1 - 0.09 - - 32 Inventive Steel 7 0.13 0.14 1.48 0.007 0.005 0.04 0.015 8 42 0.15 0.1 - - - - 0.02 - - 36 Inventive Steel 8 0.11 0.15 1.52 0.007 0.005 0.06 0.018 10 43 0.001 - 0.1 - - 0.015 0.01 - - 44 Inventive Steel 9 0.12 0.21 1.51 0.007 0.005 0.025 0.02 4 39 0.002 - - 0.1 - - 0.02 0.001 - 43 Inventive Steel 10 0.07 0.16 1.45 0.008 0.006 0.045 0.025 6 37 0.05 - 0.3 - - 0.01 0.02 - 0.01 52 Inventive Steel 11 0.11 0.21 1.49 0.008 0.003 0.047 0.019 11 44 0.01 - 0.1 - - - - 0.001 - 48 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. Conventional steels (4, 5, 6) are Japanese Patent Application Laid-open Nos. Invented steels (14, 24, 28) and conventional steels (7, 8, 9, 10) are invented steels (48, 58, 60, 61) of Japanese Patent Application Laid-Open No. 8-60292. Is invention steel F of JP-A-11-140582

Steel grade used division Primary deoxidation sequence Dissolved oxygen after the first deoxidation (ppm) Ti addition after finishing deoxidation (%) Molten steel holding time (min) Casting speed (m / min) Oxygen content after coagulation (ppm) Inventive Steel 1 Invention 1 Mn → Si 58 0.015 15 1.1 32 Invention 2 Mn → Si 55 0.015 15 1.1 34 Invention 3 Mn → Si 65 0.015 15 1.1 35 Comparative Material 1 Mn → Si 28 0.015 5 1.1 9 Comparative Material 2 Mn → Si 212 0.015 75 1.1 128 Inventive Steel 2 Invention 4 Mn → Si 67 0.053 12 1.0 54 Invention Steel 3 Invention 5 Mn → Si 61 0.016 14 1.0 32 Inventive Steel 4 Invention 6 Mn → Si 60 0.022 16 1.1 31 Inventive Steel 5 Invention 7 Mn → Si 59 0.053 15 1.0 52 Inventive Steel 6 Invention Material 8 Mn → Si 58 0.022 16 0.95 32 Inventive Steel 7 Invention Material 9 Mn → Si 62 0.016 16 0.98 36 Inventive Steel 8 Invention 10 Mn → Si 73 0.019 15 0.98 44 Inventive Steel 9 Invention 11 Mn → Si 74 0.022 15 1.0 43 Inventive Steel 10 Invention Material12 Mn → Si 69 0.026 14 1.0 52 Inventive Steel 11 Invention Material 13 Mn → Si 65 0.020 15 1.1 48 Conventional steel (1-11) is not specifically described its manufacturing conditions.

Steel grade used division Heating temperature (℃) Nitrogen atmosphere (ℓ / min) Heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount at recrystallization area (%) / cumulative pressure drop ratio (%) Cooling rate (℃ / min) Base nitrogen (ppm) Invention 2 Inventive Example 1 1250 350 170 1030 790 65/80 8 105 Inventive Example 2 1210 630 120 1040 790 65/80 8 115 Inventive Example 3 1150 780 110 1040 790 65/80 8 120 Comparative Example 1 1050 18 50 1040 790 65/80 8 68 Comparative Example 2 1300 950 240 1040 790 65/80 8 224 Invention 1 Inventive Example 4 1210 610 130 1020 800 65/80 7 114 Invention 3 Inventive Example 5 1220 600 130 1040 810 65/80 8 118 Comparative Material 1 Comparative Example 3 1220 610 120 1030 790 65/80 8 110 Comparative Material 2 Comparative Example 4 1210 600 130 1030 790 65/80 7 120 Invention 4 Inventive Example 6 1190 680 140 1020 800 65/80 8 275 Invention 5 Inventive Example 7 1240 550 110 1010 810 60/75 6 112 Invention 6 Inventive Example 8 1180 460 170 1030 800 60/75 6 80 Invention 7 Inventive Example 9 1210 740 110 1020 800 60/75 7 300 Invention Material 8 Inventive Example 10 1220 620 100 1000 790 60/75 7 100 Invention Material 9 Inventive Example 11 1210 620 110 1030 780 60/75 7 115 Invention 10 Inventive Example 12 1180 550 150 1020 790 60/75 7 120 Invention 11 Inventive Example 13 1210 580 120 1040 800 60/70 7 90 Invention Material12 Inventive Example 14 1190 640 130 1020 800 60/70 7 100 Invention Material 13 Inventive Example 15 1210 590 120 1020 790 65/75 8 132 Conventional Materials 11 1200 120 1050 850 3 89 Cooling of the invention is performed at a cooling rate of 5 ° C./sec up to 600 ° C. after the completion of ferrite transformation, and then air-cooled thereafter. , Conventional steels (1-10) do not have specific hot rolling conditions.

Alloying element ratio after nitriding treatment to show the effect of the present invention Ti / O Ti / N N / B Al / N V / N (Ti + 2Al + 4B + V) / N Inventive Example 1 4.1 1.3 15 3.8 1.0 10.2 Inventive Example 2 4.1 1.2 16.4 3.5 0.9 9.3 Inventive Example 3 4.1 1.2 17.1 3.3 0.8 8.9 Comparative Example 1 4.1 2.1 9.7 5.9 1.5 15.7 Comparative Example 2 4.1 0.6 32 1.8 0.5 4.8 Inventive Example 4 4.4 1.2 16.3 3.6 0.9 9.4 Inventive Example 5 4.0 1.2 16.9 3.4 0.8 9.1 Comparative Example 3 15.6 1.3 15.7 3.6 0.9 9.7 Comparative Example 4 1.1 1.2 17.1 3.3 1.0 10.2 Inventive Example 6 9.3 1.8 27.5 2.5 0.4 7.4 Inventive Example 7 4.7 1.3 37.3 5.4 1.8 14.0 Inventive Example 8 6.5 2.5 16.0 2.5 6.3 14.0 Inventive Example 9 9.6 1.7 20.0 3.0 1.7 9.5 Inventive Example 10 6.3 2.0 10.0 2.5 9.0 16.4 Inventive Example 11 4.2 1.3 14.4 3.5 1.7 10.3 Inventive Example 12 4.1 1.5 12.0 5.0 0.8 12.7 Inventive Example 13 4.7 2.2 22.5 2.8 2.2 10.2 Inventive Example 14 4.8 2.5 16.7 4.5 2.0 13.7 Inventive Example 15 4.0 1.4 12 3.6 - 8.9 Conventional Steel 1 0.4 4.1 13.8 0.6 - 5.7 Conventional Steel 2 3.8 2.5 96.0 0.8 - 4.0 Conventional Steel 3 0.7 0.8 105.8 0.4 - 1.5 Conventional Steel 4 5.2 4.1 4.0 0.8 8.8 15.5 Conventional Steel 5 4.8 6.5 4.0 1.1 18.5 28.1 Conventional Steel 6 0.4 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

division Precipitate properties Oxide properties Base material texture characteristics Count (pcs / mm ² ) Average size (㎛) Average interval (㎛) Count (pcs / mm ² ) Average size (㎛) Thickness (mm) Yield strength (MPa) Tensile Strength (MPa) Elongation (%) FGS (μm) Ferrite Percentage (%) Impact toughness of -40 ℃ (J) Inventive Example 1 2.3 X ^{10 8} 0.017 0.24 2.3 X 10 ² 1.8 25 396 553 38 11 86 364 Inventive Example 2 3.3 X ^{10 8} 0.016 0.25 2.1X10 ² 1.4 25 395 551 39 9 88 365 Inventive Example 3 2.3 X ^{10 8} 0.015 0.27 2.3 X 10 ² 1.2 25 396 550 39 10 87 365 Comparative Example 1 4.3X10 ⁶ 0.153 1.6 3.5X10 ¹ 2.2 25 393 554 26 14 75 225 Comparative Example 2 5.4 X 10 ⁶ 0.148 1.8 2.4 X 10 ¹ 2.3 25 392 550 27 15 72 220 Inventive Example 4 3.4 X ^{10 8} 0.024 0.32 3.2 X 10 ² 1.2 30 396 558 38 11 82 345 Inventive Example 5 2.6 X ^{10 8} 0.018 0.36 2.6 X 10 ² 1.2 30 396 562 38 10 83 364 Comparative Example 3 2.2 X 10 ⁶ 0.165 1.3 3.3X10 ¹ 2.6 30 386 542 25 13 70 224 Comparative Example 4 3.4 X 10 ⁶ 0.148 1.2 3.2 X 10 ¹ 2.8 30 394 562 24 15 71 202 Inventive Example 6 3.2 X ^{10 8} 0.026 0.35 1.8 X 10 ² 1.3 30 390 563 38 10 83 342 Inventive Example 7 3.4 X ^{10 8} 0.024 0.42 3.2 X 10 ² 1.2 35 390 564 39 10 85 364 Inventive Example 8 4.3X10 ⁸ 0.014 0.35 2.1X10 ² 1.6 35 392 542 36 11 83 351 Inventive Example 9 5.4 X ^{10 8} 0.028 0.29 2.4 X 10 ² 1.5 35 391 536 37 10 84 352 Inventive Example 10 5.2 X ^{10 8} 0.021 0.44 2.3 X 10 ² 1.5 35 394 566 36 11 84 365 Inventive Example 11 3.5 X ^{10 8} 0.029 0.35 2.7 X 10 ² 1.7 40 390 566 37 12 86 349 Inventive Example 12 2.5 X ^{10 8} 0.025 0.26 1.7 X 10 ² 1.5 40 396 542 38 11 85 365 Inventive Example 13 3.7 X ^{10 8} 0.026 0.42 2.6 X 10 ² 1.5 40 402 565 38 12 84 345 Inventive Example 14 2.3 X ^{10 8} 0.025 0.47 1.8 X 10 ² 1.4 40 412 562 39 11 85 346 Inventive Example 15 3.6X10 ⁸ 0.023 0.31 1.8 X 10 ² 1.6 30 398 548 39 10 86 357 Conventional Steel 1 35 406 436 - Conventional Steel 2 35 405 441 - Conventional Steel 3 25 629 681 - Conventional Steel 4 Precipitates of MgO-TiN 3.03 × 10 ⁶ pcs / mm ² 40 472 609 32 Conventional Steel 5 Precipitates of MgO-TiN 4.07 × 10 ⁶ pcs / mm ² 40 494 622 32 Conventional Steel 6 Precipitates of MgO-TiN2.80 × 10 ⁶ pcs / mm ² 50 812 912 28 Conventional Steel 7 25 629 681 - Conventional Steel 8 50 504 601 - Conventional Steel 9 60 526 648 - Conventional Steel 10 60 760 829 - Conventional Steel 11 TiN 0.2μm or less 11.1 × 10 ³ 50 401 514 18.3

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 2.4X10 ⁸ / mm ² or more, whereas in the case of conventional steel 4.07 X10 ⁶ / mm ² Since the following ranges show that the invention material has a fairly uniform and fine precipitate size compared to the conventional material, the number is also significantly increased. In addition, the present invention has a number of oxides (TiO) of about 1.7x10 ² ~ 3.26x10 ³ ranges and the average size is about 1.2-1.8㎛ range. On the other hand, in the structure of the base steel of the present invention, the steel of the present invention has a ferrite grain size (FGS) of about 9-12 μm, which is very fine compared to that of the comparative material. It consists of fractions.

As shown in Table 6, the austenitic grain size under the highest heating temperature of 1400 ° C., such as the weld heat affected zone, ranges from 52 to 64 μm in the case of the present invention, while in the prior art, it is very roughly about 180 μm or more. You can see that we go to the range. Therefore, in the present invention, it can be seen that the austenite grain suppression effect of the weld heat affected zone during welding is very excellent. In addition, the ferrite phase fraction of the steel of the present invention at a heat input of 100 kJ / cm is about 70% or more.

As described above, the present invention can further improve the input heat welding heat affected zone impact toughness as it distributes stable and fine TiN precipitates even at high temperature through the immersion treatment on the low nitrogen steel slab. It is possible to provide a structural steel for welding that can secure toughness at the same time.

Claims (5)

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: 0.005% or less, B: 0.0003-0.01% , W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, low nitrogen steel slab composed of remaining Fe and other impurities and satisfying 4≤Ti / O≤10 Making step; Heating the slab at a temperature of 1100 to 1250 ° C. for 60 to 180 minutes to immerse the steel to satisfy the following relationship with Ti, B, and Al in the range of 0.008 to 0.03%; And 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14 The heated slab is hot rolled in an austenite recrystallization zone at a rolling ratio of 40% or more, and then cooled to a ferrite transformation end temperature ± 10 ° C at a rate of 1 ° C / min or more. Method of manufacturing structural steels. The steel material according to claim 1, wherein V is contained in an amount of 0.01 to 0.2%, a ratio of V and N (V / N) is 0.3 to 9, and 7≤ (Ti + 2Al + 4B + V) / N≤17. Method for producing a welded structural steel with excellent weld heat affected zone toughness characterized in that it satisfies. 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%. 1 or 2 or more selected from the group of Ca: 0.0005-0.005%, Rem: 0.005 to 0.05% of the method for producing a welded structural steel having excellent toughness of the weld heat affected zone characterized in that it contains. . The method of claim 1, wherein the steel slab is added to the deoxidation element having a greater deoxidizing power than Ti before the input of Ti to deoxidize the dissolved oxygen of molten steel to 50 to 200ppm, and then added so that Ti is 0.005 to 0.2% A method of manufacturing a welded structural steel having excellent toughness in welding heat affected zones, characterized in that it is gas treated and continuously cast. 5. The TiN precipitate according to claim 1, 2 or 4, wherein the steel is made of a composite structure of ferrite and pearlite having a microstructure of 20 μm or less, and a TiN precipitate having a size of 0.01-0.1 μm. at an interval of less than 1.0x10 ㎛ and ⁷ / ㎟ above distribution, in addition, the welding heat affected zone toughness, characterized in that the oxide of the TiO ^{0.5~2.0㎛ 1 × 10 2 -1 × 10} 3 gae / mm ² gae distribution Excellent method for manufacturing welded steels.