KR100482196B1 - Method of manufacturing steel plate to be precipitating TiO and TiN by nitriding treatment for welded structures - Google Patents

Abstract

The present invention relates to structural steels used in welded structures such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes, etc. The purpose of the present invention is to fine-tune TiO oxides while securing fine TiN precipitates by sedimentation on low-nitrogen steel slabs. The present invention provides a method for manufacturing a welded structural steel which can be uniformly distributed to improve the toughness of the weld heat affected zone.

The present invention for achieving the above object is dissolved in a tertiary deoxidation by adding a deoxidation element to molten steel to adjust the dissolved oxygen amount to 60 to 120ppm, followed by secondary deoxidation with a Ti deoxidizer and a deoxidation element having a higher deoxidizing power than Ti. After adjusting the amount of oxygen to 60 ppm or less, refinement was carried out by adding Ti so that the amount of Ti was 0.005-0.2%, and then, 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: 0.001-0.03 Casting a low nitrogen steel slab composed of%, remaining Fe and other impurities, and satisfying 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 was hot-rolled at an austenitic 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 gist of the present invention relates to a method for producing a welded structural steel having fine TiO oxides.

Description

Method for manufacturing steel plate to be precipitating TiO and TiN by nitriding treatment for welded structures} by TiN precipitate and fine TiO oxide

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 manufacturing a welded structural steel that can improve the toughness of the weld heat affected zone by finely and uniformly distributing a large amount of TiN oxide and TiO oxide 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, (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 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 a high heat input 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 subjected to the nitriding treatment, the precipitate of TiO and the TiO oxide are finely and uniformly distributed by the nitriding treatment while fundamentally blocking the occurrence of cracks on the surface of the slag of high nitrogen steel. It is an object of the present invention to provide a method for manufacturing a welded structural steel having a minimum difference in toughness of a base material and a heat affected zone in a heat input amount range.

In the method of manufacturing the welded structural steel of the present invention for achieving the above object, by adding a deoxidation element to molten steel to adjust the dissolved oxygen amount to 60 to 120ppm, the second deoxidation with Ti deoxidizer and the deoxidation element having a higher deoxidizing power than Ti. After adjusting the dissolved oxygen amount to 60ppm or less by adding tertiary deoxidation, refinement was added by adding Ti to 0.005-0.2% of Ti, followed by C: 0.03-0.17% by weight, Si: 0.01-0.5%, and 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% Casting a low nitrogen steel slab composed of O: 0.001-0.03%, remaining Fe and other impurities, and satisfying 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 low-nitrogen steels, the present invention has excellent high temperature stability through nitriding in a subsequent process. While uniformly distributing the TiN precipitates and finely dispersing the oxides of TiO, it was confirmed that the grain size of the austenite was below the threshold (about 80 µm), so that the toughness of the weld heat affected zone was not a problem.

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. When the range of 1.2 to 2.5 is controlled, the amount of solid solution Ti is extremely reduced and the high temperature stability of the TiN precipitate is increased, so that the fine TiN 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. This increases the content of nitrogen at the same Ti content, so that all the Ti atoms dissolved are easily combined with the nitrogen atom, and in a high nitrogen environment, the amount of solid solution Ti decreases, making TiN precipitates more stable at higher temperatures than in the case of low nitrogen content. It is analyzed that this is because the solubility product is lowered. In view of the fact that the present invention can promote 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, as a result of concentrated research on the oxide of TiO,

First, TiN is preferentially precipitated using a fine TiO oxide as a nucleus, and the number of TiN precipitates is increased.

Second, TiO oxide can promote the nucleation of acicular ferrite while inhibiting grain growth of gustenite, both of which have been found to be more uniformly dispersed as a fine dispersion of TiO oxide. To this end, in the present invention, Ti is divided into primary and secondary and added to obtain a finer Ti oxide.

[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 the 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 a central segregation zone in the center part during rolling, thereby creating a low temperature transformation structure in the center of the base material. .

The content of aluminum (Al) is preferably limited to 0.005-0.1%.

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 precipitation interval of the TiN and AlN precipitates, the distribution of precipitates, the frequency of complex precipitation with oxide, 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. In the present invention, in the steel slab N is managed to less than 0.005% of low probability of cast surface cracking, and then made into a high quality steel of 0.008-0.03% through the immersion treatment in the slab reheating process.

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 0.0010-0.006%.

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 0.0010%, it is not preferable because the number of Ti oxides is small, and if it exceeds 0.006%, coarse Ti oxides and other oxides such as FeO are produced.

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 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.

[Manufacturing method]

Refining Process (Deoxidation and Degassing Process)

In general, the steel refining process is first refined in a refining furnace (electric furnace, electric furnace), and then the molten steel of the refining furnace is pulled out with a ladle (LF) for secondary refining (external refining +). It is preferable to perform degassing (RH process). Normal deoxidation takes place in primary and secondary refining.

In the present invention, it is characterized by finely and uniformly dispersing the oxides of TiO by adjusting dissolved oxygen to an appropriate level in such a deoxidation process and then dividing Ti into two stages. In order to uniformly disperse the fine Ti oxide in the slab in the process of producing molten steel, it is preferable to precipitate the fine secondary oxide formed during solidification of the molten steel rather than the coarse primary oxide formed in the molten steel. Coarse primary oxide does not play a role in inhibiting austenite grain growth, and also does not play a role of nucleation of acicular ferrite in the weld heat affected zone. Therefore, the fine secondary oxide is preferably precipitated in the course of continuous slab casting of molten steel. In order to precipitate fine secondary oxides, it is important not only to control Ti concentration and oxygen concentration in molten steel but also to input Ti in two steps.

The primary Ti addition combines with the oxygen in the molten steel to form a primary oxide, which is attached to the slag of the molten steel by flotation separation, so that the oxygen concentration in the molten steel can be controlled. The secondary Ti addition is added immediately before solidification. The added secondary Ti exists as solid-solution Ti in molten steel because there is no time to bond with oxygen, and it can be uniformly dispersed because it forms fine secondary Ti oxide in the slab by combining with dissolved oxygen during solidification. Fine uniformly dispersed Ti oxide is preferable because it suppresses austenite grain growth and is advantageous in securing weldability because it can promote nucleation of acicular ferrite, which is known to have excellent toughness in the weld heat affected zone during welding.

Therefore, in the present invention, the deoxidation method for finely dispersing the TiO oxide by two-step addition of Ti, (1) dissolved oxygen control step, (2) Ti deoxidation step, (3) secondary dissolved oxygen control step, ( 4) It consists of four steps of the addition step of Ti.

(1) Dissolved Oxygen Control Stage

In the present invention, a deoxidizer is added to adjust the dissolved oxygen to 60-120 ppm. As for the deoxidizer at this time, it is more preferable to use the thing with less deoxidizing power than Ti. The deoxidizing power of the deoxidizer is as follows.

Cr <Mn <Si <Ti <Al <REM <Zr <Ca ≒ Mg

Since the dissolved oxygen amount greatly influences the production behavior of the oxide, in order to obtain the desired Ti oxide, the dissolved oxygen amount must be managed at a level of 60-120 ppm before the addition of Ti. In the present invention, the deoxygenation elements of Mn and Si are added to adjust the dissolved oxygen amount to 60-120 ppm. If the dissolved oxygen amount is less than 60 ppm, the amount of oxygen in the molten steel is too small to form an adequate amount of Ti-based oxide to show the effect of the present invention. If the dissolved oxygen is more than 120 ppm, the effect is saturated, and not only the target element but also other alloy addition elements are oxidized. This is because there is a problem of obtaining an alloy component system composed of accurate chemical components.

(2) deoxidation (primary injection) of Ti,

In the present invention, the dissolved oxygen is adjusted to 60-120 ppm, and then Ti is added to form a primary oxide of TiO. The primary oxide of TiO is attached to the slag and removed by flotation, but part of it is reduced by the deoxidizer which is added and then refined.

(3) secondary dissolved oxygen control step,

In the present invention, after adding Ti, a deoxidizing agent having a higher deoxidizing power than Ti is added to adjust the dissolved oxygen amount to 60 ppm or less. When the dissolved oxygen amount is more than 60 ppm, the Ti added in the subsequent process is present as a crystallized Ti oxide to prepare slabs. It is not preferable because it reduces the amount of fine TiN precipitates formed.

At this time, the deoxidizer is preferably Al, so that the final Al content in the molten steel is in the range of 0.005-0.1%. After Ti deoxidation, Ti oxide is partially reduced and further refined by Al injection, and the number of Ti oxides also increases. This is because the amount of dissolved oxygen decreases due to the addition of Al to control the growth of the Ti oxide, which makes the Ti oxide finer and thus makes it more difficult to float. As a result of experiment to derive the proper range of Al, if Al is less than 0.005%, the influence on reduction of Ti oxide and dissolved oxygen amount is not enough, Ti oxide coarsens and floats, and exceeds 0.1%. When the Ti oxide is completely reduced it was confirmed that the number of Ti oxide is reduced.

In the present invention, it is preferable to add Al as a tertiary deoxidation element within 10 minutes after Ti deoxidation. Ti oxide produced over time after Ti deoxidation has a problem of coagulation and flotation due to growth aggregation. For this reason, it is advantageous to inject Al immediately after Ti is uniformly mixed in molten steel and to secure a large amount of Ti oxide. However, even if Al is not added immediately after Ti deoxidation, the amount of reduction of Ti oxide is not large.

(4) Ti addition step

In the present invention, 0.005-0.2% Ti is added in the molten steel state after the third deoxidation. At this time, the addition of Ti has a different purpose from that added to Ti during secondary deoxidation. Rather than Ti oxides produced in the molten steel state, Ti oxides produced during solidification are more fine and more numerous. Therefore, when the Ti is added after tertiary deoxidation, the ratio remaining as solid solution Ti is considerably higher than that of oxides, and thus a fine amount of Ti oxide and TiN precipitates can be secured during solidification. At this time, the preferred amount of Ti is 0.005-0.2% range in consideration of the TiN, MnS, CuS, BN, VN precipitation behavior effective in improving the toughness of the weld heat affected zone as well as the appropriate Ti oxide distribution for showing the effect of the present invention.

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.

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, Osval A method of maximally suppressing Ostwald ripening was derived. In addition to the fundamental prevention of slab cracking problems commonly found in high-nitrogen steels, the sedimentation effects in slab furnaces include 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, it is preferable to heat 60-180 minutes at 1100-1250 ° C as described above, and then hot roll at a rolling ratio of 40% or more at the austenite recrystallization 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 leads to grain growth of the recrystallized fine ferrite, it is difficult to secure the ferrite grain size of the base material 20㎛ or less.

· 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 used as a sample, dissolved in a converter, manufactured into slabs by the continuous casting method, and hot rolled materials were manufactured under the conditions of Tables 2 and 3. The composition ratio between the alloying element elements of this hot-rolled material is shown in Table 4.

The test pieces for evaluating the mechanical properties of the base metal from the hot rolled plates as described above were taken at the center of the plate thickness of the rolled material, the tensile test piece was taken in the rolling direction, and the Charpy impact 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 - - 19 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 - - 50 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 - - 21 Inventive Steel 4 0.11 0.13 1.49 0.006 0.005 0.02 0.02 5 40 0.001 - - - - - 0.05 - - 23 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 - 50 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 - - 22 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 - - 21 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 - - 20 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 - 21 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 27 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 - 24 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 Molten oxygen (ppm) Primary Ti addition amount (%) Al addition amount (%) Dissolved oxygen after Al addition (ppm) 2nd Ti addition amount (%) Molten steel holding time (min) Oxygen content after coagulation (ppm) Casting condition Casting speed (m / min) Specific quantity (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 63 0.008 0.043 31 0.010 15 18 1.0 0.35 Invention 2 Mn → Si 60 0.008 0.045 28 0.010 15 19 1.0 0.35 Invention 3 Mn → Si 65 0.008 0.042 27 0.010 15 18 1.0 0.35 Comparative Material 1 Mn → Si 26 0.008 0.045 10 0.010 15 8 1.0 0.35 Comparative Material 2 Mn → Si 206 0.008 0.045 184 0.010 65 124 1.0 0.35 Inventive Steel 2 Invention 4 Mn → Si 68 0.015 0.07 50 0.033 16 50 0.95 0.32 Invention Steel 3 Invention 5 Mn → Si 67 0.009 0.06 38 0.014 15 21 1.0 0.32 Inventive Steel 4 Invention 6 Mn → Si 60 0.01 0.02 36 0.020 14 23 1.0 0.32 Inventive Steel 5 Invention 7 Mn → Si 69 0.023 0.09 50 0.040 15 50 1.1 0.35 Inventive Steel 6 Invention Material 8 Mn → Si 65 0.01 0.025 34 0.018 15 22 1.0 0.35 Inventive Steel 7 Invention 9 Mn → Si 61 0.009 0.04 30 0.010 16 21 1.1 0.35 Inventive Steel 8 Invention 10 Mn → Si 72 0.008 0.06 32 0.011 18 20 1.0 0.35 Inventive Steel 9 Invention 11 Mn → Si 67 0.008 0.025 30 0.014 14 21 0.98 0.32 Inventive Steel 10 Invention Material12 Mn → Si 75 0.007 0.045 39 0.020 11 27 1.0 0.32 Inventive Steel 11 Invention Material 13 Mn → Si 68 0.006 0.047 32 0.018 12 24 1.1 0.35 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 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 The cooling of the invention controls the cooling rate to 600 ° C., which is the temperature after the ferrite transformation is completed, and then air-cooled thereafter. 1-10) does not specifically show the 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 7.4 1.3 15 3.8 1.0 10.2 Inventive Example 2 7.4 1.2 16.4 3.5 0.9 9.3 Inventive Example 3 7.4 1.2 17.1 3.3 0.8 8.9 Comparative Example 1 7.4 2.1 9.7 5.9 1.5 15.7 Comparative Example 2 7.4 0.6 32 1.8 0.5 4.8 Inventive Example 4 7.8 1.2 16.3 3.6 0.9 9.4 Inventive Example 5 7.8 1.2 16.9 3.4 0.8 9.1 Comparative Example 3 17.5 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 10.0 1.8 27.5 2.5 0.4 7.4 Inventive Example 7 7.5 1.3 37.3 5.4 1.8 14.0 Inventive Example 8 8.7 2.5 16.0 2.5 6.3 14.0 Inventive Example 9 10.0 1.7 20.0 3.0 1.7 9.5 Inventive Example 10 9.1 2.0 10.0 2.5 9.0 16.4 Inventive Example 11 8.8 1.3 14.4 3.5 1.7 10.3 Inventive Example 12 9.0 1.5 12.0 5.0 0.8 12.7 Inventive Example 13 9.5 2.2 22.5 2.8 2.2 10.2 Inventive Example 14 9.3 2.5 16.7 4.5 2.0 13.7 Inventive Example 15 7.9 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 2.4 X ^{10 8} 0.017 0.22 2.4 X 10 ² 1.7 25 396 553 38 11 86 372 Inventive Example 2 3.5 X ^{10 8} 0.016 0.20 2.5 X 10 ² 1.4 25 395 551 39 9 88 375 Inventive Example 3 2.7 X ^{10 8} 0.015 0.21 2.4 X 10 ² 1.2 25 396 550 39 10 87 366 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 3.5 X ^{10 8} 0.024 0.31 3.3 X 10 ² 1.2 30 396 558 38 11 82 347 Inventive Example 5 2.7 X ^{10 8} 0.018 0.30 2.7 X 10 ² 1.2 30 396 562 38 10 83 366 Comparative Example 3 2.2 X 10 ⁶ 0.165 1.3 3.4X10 ¹ 2.6 30 386 542 25 13 70 224 Comparative Example 4 3.4 X 10 ⁶ 0.148 1.2 3.3 X 10 ¹ 2.8 30 394 562 24 15 71 202 Inventive Example 6 3.3 X ^{10 8} 0.016 0.34 2.5 X 10 ² 1.3 30 390 563 38 10 83 346 Inventive Example 7 3.5 X ^{10 8} 0.014 0.36 3.3 X 10 ² 1.2 35 390 564 39 10 85 365 Inventive Example 8 4.4 X ^{10 8} 0.015 0.34 2.2 X 10 ² 1.6 35 392 542 36 11 83 355 Inventive Example 9 5.5X10 ⁸ 0.018 0.28 2.5 X 10 ² 1.5 35 391 536 37 10 84 352 Inventive Example 10 5.6 X ^{10 8} 0.021 0.34 2.3 X 10 ² 1.5 35 394 566 36 11 84 366 Inventive Example 11 3.7 X ^{10 8} 0.022 0.32 2.7 X 10 ² 1.7 40 390 566 37 12 86 354 Inventive Example 12 2.6 X ^{10 8} 0.021 0.24 2.2 X 10 ² 1.5 40 396 542 38 11 85 367 Inventive Example 13 3.8X10 ⁸ 0.024 0.35 2.6 X 10 ² 1.5 40 402 565 38 12 84 356 Inventive Example 14 2.3 X ^{10 8} 0.022 0.37 2.1X10 ² 1.4 40 412 562 39 11 85 364 Inventive Example 15 3.7 X ^{10 8} 0.020 0.30 1.9 X 10 ² 1.6 30 398 548 39 10 86 362 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.

division Austenitic grain size of welding heat affected zone (㎛) Microstructure with welding heat effect of 100kJ / cm heat input Reproduction Weld Heat Affected Zone -40 ℃ Impact Toughness (J) (Maximum Heating Temperature: 1400 ℃) 1200 (℃) 1300 (℃) 1400 (℃) Ferrite Percentage (%) Ferrite Average Grain Size (㎛) Δt _800-500 = 60 seconds Δt _800-500 = 120 seconds Δt _800-500 = 180 seconds Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Inventive Example 1 20 34 52 75 15 373 -74 337 -67 295 -63 Inventive Example 2 22 35 51 78 13 385 -76 352 -69 307 -64 Inventive Example 3 19 35 52 75 13 368 -72 333 -67 296 -63 Comparative Example 1 54 86 184 38 24 123 -43 43 -34 28 -28 Comparative Example 2 65 92 196 36 26 101 -40 30 -32 17 -25 Inventive Example 4 22 38 60 76 14 354 -71 330 -68 289 -65 Inventive Example 5 24 41 52 78 15 367 -71 336 -67 299 -62 Comparative Example 3 56 80 187 40 26 103 -39 56 -32 24 -24 Comparative Example 4 63 88 193 39 28 66 -28 39 -30 10 -21 Inventive Example 6 19 32 48 75 14 384 -73 357 -69 313 -63 Inventive Example 7 21 35 50 77 14 366 -71 341 -67 298 -63 Inventive Example 8 27 37 51 74 13 366 -71 342 -67 299 -62 Inventive Example 9 24 36 50 78 15 369 -72 333 -67 285 -63 Inventive Example 10 22 34 51 75 14 384 -72 346 -66 294 -63 Inventive Example 11 26 35 49 75 14 356 -71 329 -68 286 -68 Inventive Example 12 27 39 58 74 15 354 -71 323 -67 278 -62 Inventive Example 13 23 38 55 74 14 355 -71 323 -67 255 -62 Inventive Example 14 25 35 51 70 15 344 -71 328 -67 249 -63 Inventive Example 15 23 36 48 76 16 354 -72 334 -68 295 -94 Conventional Steel 1 -58 Conventional Steel 2 -55 Conventional Steel 3 -54 Conventional Steel 4 230 93 132 (0 ℃) Conventional Steel 5 180 87 129 (0 ℃) Conventional Steel 6 250 47 60 (0 degrees Celsius) Conventional Steel 7 -60 -61 Conventional Steel 8 -59 -48 Conventional Steel 9 -54 -42 Conventional Steel 10 -57 -45 Conventional Steel 11 219 (0 ℃)

As shown in Table 6, the austenitic grain size at the maximum heating temperature of 1400 ℃ conditions, such as the welding heat affected zone has a range of 48-60㎛ in the case of the present invention, while in the prior art is very roughly about 180㎛ 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 impact toughness of the heat input welding heat affected zone by further distributing stable and fine TiN precipitates even at high temperature through not only the fine Ti oxide but also the sedimentation treatment on the low nitrogen steel slab. It is possible to provide a structural steel for welding that can at the same time secure the toughness of the heat welding heat affected zone.

Claims (5)

After adding deoxidation element to molten steel to adjust dissolved oxygen amount to 60 ~ 120ppm, secondary deoxidation with Ti deoxidizer and adding deoxidation element with more deoxidizing power than Ti, and adjusting dissolved oxygen amount to 60ppm or less. Ti was refined so that the added amount was 0.005-0.2%, followed by C: 0.03-0.17% by weight, 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: 0.001-0.03%, and the remaining Fe and other impurities Casting the low nitrogen steel slab satisfying 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 was 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 for producing welded steel having fine TiO oxide. 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 is satisfied. A method for producing a welded structural steel having TiN precipitates and fine TiO oxides by immersion treatment. 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%. At least one selected from the group consisting of one or two or more selected from the group of Ca: 0.0005-0.005% and REM: 0.005 to 0.05%, and the REM is one or two of Ce, La, Y and Hf. A process for producing a welded structural steel having a TiN precipitate and a fine TiO oxide by an immersion treatment, characterized by being a phase. The Ti denitrate and fine Ti oxide according to claim 1, wherein the tertiary deoxidation is performed by adding Al within 10 minutes after the second deoxidation with Ti so that the final content of Al is 0.005-0.1%. Method for producing a welded structural steel having a. According to claim 1, wherein the steel is made of a microstructure of the composite structure of ferrite and perlite of 20㎛ or less, TiN precipitates of 0.01-0.1㎛ distributed more than 1.0x10 ⁷ / mm2 at intervals of 0.5㎛ or less And 0.5-2.0 μm of TiO oxide having a size of 1 × 10 ² -1 × 10 ³ / mm ^{2. The} method of manufacturing a welded structural steel having TiN precipitate and fine TiO oxide by immersion treatment.