KR100435489B1 - method for manufacturing high strength steel plate to be precipitating TiN and ZrN for welded structures - 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 secure fine TiN precipitates and ZrN precipitates by immersion treatment in low nitrogen steel slabs. The base metal is made of bainite + ferrite while improving 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%, Zr: 0.001-0.03%, N: 0.008-0.030%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.003-0.05%, O: 0.01% or less Making a low nitrogen steel slab composed of other impurities,

Heating the slab at a temperature of 1100-1250 ° C. for 60-180 minutes to make N of 0.008-0.03%, while N and Ti, Zr, B, and Al are immersed to satisfy the following relationship,

1.2≤Ti / N≤2.5, 0.3≤Zr / N≤2.0, 10≤N / B≤40, 2.5≤Al / N≤7, 6.8≤ (Ti + Zr + 2Al + 4B) / N≤17,

The heated slab is hot rolled in an austenite recrystallization zone with a reduction ratio of 40% or more, and then cooled at a rate of 5-20 ° C / sec to a bainite transformation end temperature ± 10 ° C. The technical gist of the present invention relates to a method for producing a high strength welded structural steel having TiN and ZrN precipitates.

Description

Method for manufacturing high strength steel plate to be precipitating TiN and ZrN for welded structures by sedimentation

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 fabrication of welded structural steels by improving the toughness of the base material as bainite + ferrite while improving the toughness of the weld heat affected zone by securing fine TiN precipitates and ZrN precipitates by immersing the low nitrogen steel slab. It is about a method.

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

In general, in welding, the larger the amount of heat input, the larger the amount of welding, so that the number of welding passes decreases. Therefore, it is advantageous to enable high heat input welding in consideration of welding productivity. In other words, increasing the amount of heat input in the welding will be able to widen the range of use. The range of high heat input currently used corresponds to approximately 100-200 kJ / cm, but in order to weld more thickened steel, that is, steel with a plate thickness of 50 mm or more, it is possible to have a super heat input range of 200-500 kJ / cm.

When the heat input is applied to the steel, the heat affected zone formed during welding, particularly the heat affected zone near the fusion boundary, is heated to a temperature close to the melting point by the amount of heat input. Accordingly, since the grains of the weld heat affected zone grow and coarse, and microstructures that are vulnerable to toughness such as upper pearlite 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. This steel is a microstructure of ferrite and pearlite, which has a mechanical property of 581 MPa in tensile strength and 405 MPa in yield strength.

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 above 1050 ° C. and rapidly cooled and then reheated for hot rolling. Becomes In addition, in this prior art, the surface cracks of cast steel are more likely to occur because continuous nitrogen is cast into molten steel containing N—5-0.02%. 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 may cause segregation cracks due to segregation in the uncoagulated portion.

Japanese Patent Laid-Open No. 9-194990 discloses a composite oxide composed of Al, Mn, and Si by controlling the ratio of Al and O to 0.3 ≦ Al / O ≦ 1.5 in low nitrogen steel (N ≦ 0.005%). The ferrite nucleation function in the austenite mouth in the heat affected zone and the transition temperature of the heat affected zone (maximum heating temperature of 1450 ℃, cooling cycle of 800 to 500 ℃ for 60 seconds) is -50 to -60. Welded structural steels are presented. This steel has a rather low transition temperature. In this prior art, high nitrogen steels use less than 0.005% of low nitrogen steels because they deteriorate toughness.

In addition, Japanese Patent Application Laid-Open No. 10-298708 uses MgO having a particle size of 0.005-0.1 μm as a nucleus, and has TiN around it, and has a size of 0.05-0.5 μm of MgO-TiN composite precipitate of 1 square mm. Steels containing 1.0 × 10 ^{5 to} 1.0 × 10 ⁷ sugars have been proposed. In this prior art, a steel component system is managed with 0.005-0.025% of Ti, 0.0002-0.005% of Mg, and 0.0005-0.008% of O in steel containing 0.002-0.0008% of N to obtain MgO-TiN composite precipitates. . The steel provided from this prior art improves to some extent the toughness of the weld heat affected zone of the steel by using a composite precipitate of MgO-TiN, but the impact of the weld heat affected zone 0 ° C. when a high heat input welding of about 200 kJ / cm is applied. Toughness is 130J, toughness is not very good.

To date, many techniques for improving the toughness of weld heat affected zones in high heat input welding have been known, but the problem of weld heat affected zones during superheated welding maintained at 1350 ℃ for a long time while basically solving the problem of cast surface cracks in high nitrogen steel. There have been no reports of significant improvements in toughness and base metal strength. In particular, there are few technologies in which the toughness of the weld heat affected zone shows the same level as that of the base metal.

The present invention is to minimize the difference in the toughness of the base material and the weld heat affected zone by distributing TiN precipitates and ZrN precipitates while fundamentally blocking the generation of slag surface cracks of high nitrogen steel through the immersion treatment on the low nitrogen steel slab, The present invention provides a method for manufacturing a welded structural steel in which an appropriate fraction of bainite is secured to increase the base material strength.

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: 0.008-0.030%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.003-0.05%, O: 0.01% or less Making a low nitrogen steel slab composed of other impurities,

Heating the slab at a temperature of 1100-1250 ° C. for 60-180 minutes to make N of 0.008-0.03%, while N and Ti, Zr, B, and Al are immersed to satisfy the following relationship,

1.2≤Ti / N≤2.5, 0.3≤Zr / N≤2.0, 10≤N / B≤40, 2.5≤Al / N≤7, 6.8≤ (Ti + Zr + 2Al + 4B) / N≤17,

And hot-rolling the heated slab at a reduction rate of 40% or more in the austenite recrystallization zone, and then cooling it at a rate of 5-20 ° C / sec to a bainite transformation end temperature ± 10 ° C.

Hereinafter, the present invention will be described in detail.

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

The present inventors have studied ways to improve the toughness of the weld heat affected zone while preventing surface cracks that may occur in high-nitrogen steels, and to improve the strength of the steel (base metal), instead of making steel slabs from low-nitrogen steels. In the subsequent process, TiN precipitates and ZrN precipitates having excellent high temperature stability are distributed through the nitriding, so that the toughness of the weld heat affected zone is not a problem if the grain size of the austenite is controlled below the threshold (about 80㎛). When accelerated cooling of the rolled material during the rolling process after the treatment, it was confirmed that the structure of bainite can be easily controlled to improve the strength of the base material.

Starting from this point of view, the present inventors have been able to derive the following method for improving the base material strength while improving the toughness of the weld heat affected zone through the immersion treatment.

[1] dispersing TiN precipitates and ZrN precipitates by immersion in low nitrogen steel slabs,

[2] By controlling the initial ferrite grain size of the steel to less than the critical level, the former austenite of the weld heat affected zone is controlled to about 80 μm or less. Also,

[3] By reducing the Ti / N ratio, ZrN, BN, and AlN precipitates are effectively precipitated to increase the ferrite fraction in the weld heat affected zone. In particular, the ferrite shape is induced into needles or polygons effective for toughness improvement. Improve the toughness of the weld heat affected zone. Meanwhile,

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

[1] TiN precipitates

The present inventors have focused on the fact that when the heat input welding is applied to the base material in which the TiN precipitates are distributed, the cause of the TiN precipitate losing the inhibitory effect of the austenite grain growth is caused by the diffusion of solid Ti atoms decomposed by the welding heat. As a result of examining the characteristics of TiN precipitates according to the / N ratio, we found a new fact that in high nitrogen environment (low Ti / N ratio), the concentration of solid solution Ti and diffusion of atomic Ti atoms are reduced and the high temperature stability of TiN precipitate is improved. It became. More interestingly, even if steel slabs are made of low-nitrogen steel of 0.005% or less, which is less prone to cast surface cracking, and then made into high-nitrogen steel by immersion treatment in slab heating furnace during the rolling process, the ratio of Ti / N is 1.2 to 2.5. In this case, the amount of solid solution Ti was extremely reduced and the high temperature stability of TiN precipitates was increased, resulting in the distribution of fine TiN precipitates of 0.01-0.1㎛ size over 1.0x10 ⁷ / mm2 at intervals of 0.5㎛ or less. .

As described above, when the ratio of Ti / N is controlled to 2.5 or less (increasing the content of N) according to the present invention, the solubility area showing the high temperature stability of TiN is also lowered. Increasing the nitrogen content at the same Ti content combines all the dissolved Ti atoms with the nitrogen atom, increasing the fine TiN precipitation amount, so that the solubility product showing the stability of the precipitate at high temperature such as the weld heat affected zone is reduced. Therefore, in a high nitrogen environment, TiN precipitates are more stable than nitrogen-containing precipitates because the amount of solid solution Ti decreases. At this time, the important thing is that the presence of solid N due to high nitrogen can promote aging, so the ratio of Zr / N, N / B, Al / N, V / N and these are collectively managed to manage N as ZrN, BN To be precipitated with AlN, VN. When controlling the ratio of Zr / N with the ratio of Ti / N by immersing the low nitrogen steel slab according to the present invention, it can be found that a large amount of fine TiN precipitates and ZrN are precipitated. This ZrN precipitate is stable even at high temperatures, and is therefore effective in suppressing the growth of the austenite.

[2] ferrite grain size management

According to the study of the present invention, it is important to make the size of the ferrite to 20 μm or less while the microstructure of the base material is a composite structure of ferrite + bainite in order to make the size of the old austenite to 80 μm. At this time, the refinement of the ferrite may be obtained by controlling the growth of the ferrite grains generated in the cooling process after hot rolling as well as the austenite grain refinement by the steel working during hot rolling. For this purpose, it was confirmed that it is very effective to properly deposit and distribute carbides (VC, WC) effective for ferrite grain growth.

[3] microstructure, welded heat affected zones

The facts of the present invention revealed that the toughness of the weld heat affected zone was not only the size of the austenite grains but also the amount of ferrite precipitated at the former austenite grain boundaries when the heat input welding was applied to the base material and heated to 1400 ° C. or more. Size and shape also influence. In particular, it is important to induce the transformation of polygonal ferrite and acicular ferrite in the austenite mouth. In the present invention, ZrN, AlN, BN precipitates are used for this purpose.

[4] bainite tissue fraction control

The present inventors can easily control the fraction of bainite, which can improve the strength of the base material when the accelerated cooling rate is controlled after being subjected to sedimentation and hot rolling in the slab reheating process. It was confirmed that even with + ferrite, the properties of the weld heat affected zone did not adversely affect.

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 pearlite, martensite and degenerate pearlite, are transformed during cooling to lower the low temperature impact toughness of structural steel, and also reduce the hardness of the welded steel or Increasing strength results 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 effective in improving deoxidation, weldability, hot workability and strength in steel, and precipitates in the form of MnS around Ti-based oxides, which affects the formation of needle-shaped and polygonal ferrites, which is effective for improving the toughness of weld heat affected zones. Crazy Such Mn forms a solid solution to form a solid solution in the matrix structure to enhance the solid solution to secure the strength and toughness, for this purpose it is preferably contained 0.4% or more. However, in case of exceeding 2.0%, tissue heterogeneity caused by Mn segregation not only has a detrimental effect on the toughness of weld heat affected zone, but also causes macro segregation and micro segregation depending on the segregation mechanism during steel solidification. It promotes the formation of central segregation zone, which acts as a cause of creating low temperature transformation tissue in the center of the base material.

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

Al is an element that helps Ti form fine TiN precipitates by forming fine AlN precipitates in steel as well as an element necessary as a deoxidizer, and forming oxygen and Al oxides to prevent Ti from reacting with oxygen. For this purpose, it is preferable to add more than 0.0005% of Al, but if it exceeds 0.1%, AlN is precipitated and the formation of Weidmanstatten ferrite and phase martensite, in which the remaining solid solution Al is vulnerable to toughness during cooling of the weld heat affected zone. To reduce the toughness of the high heat input welding heat affected zone.

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

Ti is indispensable in the present invention because it combines with N to form fine TiN precipitates that are stable at high temperatures. 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 zirconium (Zr) is preferably limited to 0.001-0.03%.

Zr is an indispensable element in the present invention because it combines with N to form a fine ZrN precipitate that is stable at high temperatures. In order to obtain such a fine Zr precipitation effect, it is preferable to add Zr 0.001% or more. However, if it exceeds 0.03%, coarse ZrN precipitates and Zr oxides are formed in molten steel, which is not preferable because it adversely affects the toughness of the base metal and the weld heat affected zone.

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 and low temperature cracking of the weld heat affected zone.

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

N is an indispensable element for forming TiN, ZrN, AlN, BN, VN, NbN, etc., and maximally suppresses the growth of old austenite grains in weld heat affected zones during high heat input welding. , Increase the amount of precipitates, such as NbN. In particular, since the TiN and ZrN precipitates have a remarkable effect on the size, precipitate interval, precipitate distribution, complex precipitation frequency with oxide, and high temperature stability of the precipitate itself, the content is preferably set at 0.008% or more. 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, N is managed in the steel slab to 0.005% or less, which is less likely to cause surface cracks, and then, the slab reheating process is made of high nitrogen steel of 0.008-0.03%.

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

Tungsten is an element that uniformly precipitates in the base material as tungsten carbide (WC) after hot rolling, effectively inhibiting ferrite grain growth after ferrite transformation, and also suppressing the growth of the initial austenite grains in the heating zone 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) 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.

The content of sulfur (S) is preferably 0.003-0.05%.

S is an element that improves the strength of the weld heat affected zone, and reacts with Cu to form CuS to improve strength (or hardness), and complex precipitation on TiN precipitate to improve high temperature stability of TiN precipitate. For this purpose, S should be contained 0.003% or more, but if it is more than 0.05%, the effect is saturated, and there is a fear that the slab may be cracked under the surface of the slab at the time of playing, and a low melting point compound such as FeS may be formed during welding. It is not preferable because it may promote welding hot cracking.

The content of oxygen (O) is preferably limited to 0.01% or less.

When oxygen exceeds 0.01%, Ti element is not preferable because Ti element is formed of Ti oxide in molten steel and TiN precipitate cannot be formed. Also, inclusions such as coarse Fe oxide and Al oxide are formed to adversely affect the toughness of the base metal. It is not desirable because it is crazy.

In the present invention, the ratio of Ti / N of steel is 1.2-2.5, the ratio of Zr / N is 0.3-2.0, the ratio of N / B is 10-40, the ratio of Al / N is 2.5-7, (Ti + Zr + 2Al + 4B) Nitrogen is preferably soaked so that the ratio / N satisfies 6.8 to 17.

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 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 determined and no uniform distribution of TiN is obtained. Also, excess Ti remaining without precipitation into TiN is in solid solution, which is bad for the toughness of the weld heat affected zone. Affect If 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 weld heat-oriented part.

The ratio of Zr / N is preferably 0.3 to 2.0.

In the present invention, when the Zr / N ratio is less than 0.3, the amount of precipitation of ZrN that prevents grain growth of the weld heat affected zone during welding is insufficient, and when the Zr / N ratio exceeds 2.0, the effect is saturated to reduce the toughness of the weld heat affected zone. It is not preferable because it is made.

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.

The ratio of (Ti + Zr + 2Al + 4B) / N is preferably 6.8 to 17.

In the present invention, when the ratio of (Ti + Zr + 2Al + 4B) / N is less than 6.8, the growth inhibition of the austenite grain growth of the weld heat affected zone, the generation of fine polygonal ferrite at the grain boundary, the amount of solid solution nitrogen, the needle shape in the grain and Insufficient size and number of distribution of TiN, ZrN, AlN, BN, and VN precipitates for the formation of polygonal ferrites and control of the fraction of tissues, and the effect is not effective when (Ti + Zr + 2Al + 4B) / N exceeds 17 Is saturated. If V is added, the ratio of (Ti + Zr + 2Al + 4B + V) / N is preferably 7-19.

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 at the TiN precipitate boundary for improving the toughness of the weld heat affected zone. When the V / N ratio exceeds 9, the size of the VN precipitates deposited on the TiN precipitate boundary is coarsened, so that the VN precipitation frequency deposited on the TiN precipitate boundary is reduced, thereby reducing the ferrite phase fraction effective for the toughness of the weld heat affected zone. Let's do it.

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 that improves the strength of the weld heat affected zone. CuS precipitates and solid solution strengthening effects are not sufficient at less than 0.1%, and when it exceeds 1.5%, the effect is saturated and rather hardenability of the weld heat affected zone. It is not preferable because it increases the toughness and decreases the toughness and dilutes the weld metal during welding, thereby decreasing the toughness of the weld metal.

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 at least 0.01%. However, Nb is not preferable 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 impair the cleanliness of the steel. As REM, 1 type, or 2 or more types, such as Ce, La, Y, and Hf, may be used, and any of the above effects can be obtained.

[Method of manufacturing welded structural steel]

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.

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

Cr <Mn <Si <Ti <Al <REM <Zr <Ca ≒ Mg

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

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

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

On the other hand, according to the 'Stoke Law' widely used in steelmaking, the higher the density of inclusions, the more difficult the inclusions are. Injuries are difficult Accordingly, since inclusions increase in the steel, the addition of a deoxidation element forming a high density inclusion may be utilized as an additional advantage according to the distribution of oxides. It does not affect the effect of.

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

Slab manufacturing process

Since molten steel of the present invention is a low nitrogen steel, the casting speed may be either high speed or low speed during continuous casting. In order to obtain a good 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 through the nitriding treatment in the slab 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, Oswald life ( Ostwald ripening).

In addition to the fact that the nitriding effect in slab furnaces can make slabs from low-nitrogen molten steel, it is possible to fundamentally prevent the problem of slab surface cracks commonly encountered in high-nitrogen steels. have. 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 furnace to increase the amount of fine TiN precipitates.

On the other hand, in the present invention, it is preferable that the nitriding treatment is performed while heating the slab at 1100-1250 ° C. for 60-180 minutes, so that the nitrogen concentration of the slab is 0.008-0.03%. First, it is preferable to set the amount of nitrogen in the slab to 0.008-0.03%. In order to secure an appropriate amount of TiN precipitation in the slab, nitrogen must be contained in 0.008% or more, but when it exceeds 0.03%, it diffuses into the slab. This is because the amount of nitrogen deposited on the surface of the slab increases more than the amount of nitrogen precipitated with fine TiN, which causes hardening on the surface of the slab, which affects the subsequent rolling process.

In addition, the slab heating temperature is set to 1100-1250 ° C. The reason is that when the heating temperature is less than 1100 ° C, the number of fine TiN precipitates is reduced because of the small driving force to diffuse the precipitated nitrogen, and the number of TIN precipitates is also increased. This is because there is a problem that the manufacturing cost is increased because the heating time must be increased to make. 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 exhibited, and when the slab heating time is longer than 180 minutes, not only does the cost of the unworking industry increase but also the austenite grain growth in the slab causes subsequent rolling process. It is not desirable because it affects.

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 Zr / N is 0.3 to 2.0, the ratio of N / B is 10 to 40, and the ratio of Al / N is 2.5. It is preferable to immerse N so that the ratio of -7 and V / N may be 0.3-9, and the ratio of (Ti + Zr + 2Al + 4B + V) / N becomes 7-19.

Hot rolling process

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

In hot rolling, the austenite grain size is affected by the temperature, time in the reheating furnace, and the rolling amount. The grain size of the austenite affects the quenchability, so that the desired bainite fraction can be obtained by controlling it. In order to increase the bainite fraction, it is recommended to set the austenite grain size to 20 µm or more. If the austenite grain size becomes larger than 70 µm, the hardenability becomes too large during transformation, which is likely to cause martensite transformation. .

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

· Microstructure of steel

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

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

Distribution of precipitates (TiN precipitates)

In the base material of the present invention, it is preferable that more than 1.0 × 10 ⁷ TiN precipitates are distributed per 1 mm ^{2 in} a size of 0.01-0.1 μm. If the precipitate size is less than 0.01㎛, it is easily re-used in the base metal during the high heat input welding, so that the effect of inhibiting austenite grain growth is insufficient, and if it exceeds 0.1㎛, the pinning (grain growth inhibition) effect on the austenite grain is small. It behaves like highly coarse nonmetallic inclusions and has a detrimental effect on mechanical properties. If the number of these fine precipitates is less than 1.0 × 10 ⁷ per 1 mm ² , it is difficult to control less than 80 μm, which is the critical austenite grain size of the weld heat-affected zone when welding more than a large heat input. Since these precipitates are more advantageous in suppressing the Ostwald ripening phenomenon in which the precipitates are coarse when uniformly distributed, it is preferable to control the interval of the TiN precipitates to 0.5 μm or less.

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

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 by melting them in a converter, and then hot-rolled sheets were manufactured under the conditions of Tables 2 and 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 Zr Cu Ni Cr Mo Nb V Ca REM O (ppm) Inventive Steel 1 0.12 0.13 1.54 0.006 0.005 0.04 0.014 7 40 0.005 0.01 0.1 - - - - 0.01 - - 11 Inventive Steel 2 0.07 0.12 1.71 0.006 0.006 0.07 0.05 10 48 0.002 0.02 - 0.2 - - - 0.01 - - 12 Invention Steel 3 0.14 0.10 1.9 0.006 0.008 0.06 0.015 3 42 0.003 0.01 - - - - - 0.02 - - 10 Inventive Steel 4 0.10 0.12 1.80 0.006 0.007 0.02 0.02 5 40 0.001 0.01 0.1 - - - - 0.05 - - 9 Inventive Steel 5 0.08 0.15 2.0 0.006 0.006 0.09 0.05 15 45 0.002 0.02 - - 0.1 - - 0.05 - - 12 Inventive Steel 6 0.10 0.14 2.0 0.007 0.005 0.025 0.02 10 47 0.004 0.01 - - - 0.1 - 0.09 - - 9 Inventive Steel 7 0.13 0.14 1.6 0.007 0.007 0.04 0.015 8 45 0.15 0.01 0.1 - - - - 0.02 - - 11 Inventive Steel 8 0.11 0.15 1.52 0.007 0.006 0.06 0.018 10 42 0.001 0.005 - - - - 0.015 0.01 - - 10 Inventive Steel 9 0.13 0.21 1.42 0.007 0.005 0.025 0.02 4 36 0.002 0.01 - - 0.1 - - 0.02 0.001 - 12 Inventive Steel 10 0.07 0.16 2.0 0.008 0.010 0.045 0.025 6 45 0.05 0.005 - 0.3 - - 0.01 0.02 - 0.01 11 Inventive Steel 11 0.11 0.21 1.48 0.007 0.006 0.047 0.019 11 48 0.01 0.005 - 0.1 - - - - - - 15 Conventional Steel 1 0.05 0.13 1.31 0.002 0.006 0.0014 0.009 1.6 22 - - - - - - - - - - 22 Conventional Steel 2 0.05 0.11 1.34 0.002 0.003 0.0036 0.012 0.5 48 - - - - - - - - - - 32 Conventional Steel 3 0.13 0.24 1.44 0.012 0.003 0.0044 0.010 1.2 127 - - 0.3 - - - 0.05 - - - 138 Conventional Steel 4 0.06 0.18 1.35 0.008 0.002 0.0027 0.013 8 32 - - - - 0.14 0.15 - 0.028 - - 25 Conventional Steel 5 0.06 0.18 0.88 0.006 0.002 0.0021 0.013 5 20 - - 0.75 0.58 0.24 0.14 0.015 0.037 - - 27 Conventional Steel 6 0.13 0.27 0.98 0.005 0.001 0.001 0.009 11 28 - - 0.35 1.15 0.53 0.49 0.001 0.045 - - 25 Conventional Steel 7 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 - - 0.3 - - - 0.036 - - - - Conventional Steel 8 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 - - 0.32 0.35 - - 0.013 - - - - Conventional Steel 9 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 - - - - 0.21 0.19 0.025 0.035 - - - Conventional Steel 10 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 - - - 1.09 0.51 0.36 0.021 0.021 - - - Conventional Steel 11 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 - - - - - - - 0.069 - - - Conventional steels (1, 2, 3) are invention steels (5, 32, 55) of Japanese Patent Application Laid-Open No. 9-194990, and conventional steels (4, 5, 6) are inventions of Japanese Patent Application Laid-Open No. 10-298708. Steels (14, 24, 28) and conventional steels (7, 8, 9, 10) are invention steels (48, 58, 60, 61) of JP-A-8-60292. Inventive Steel F of Patent Publication No. 11-140582

Steel grade used division Primary deoxidation sequence Dissolved oxygen after Al addition (ppm) Ti addition after finishing deoxidation (%) Casting speed (m / min) Specific quantity (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 19 0.014 1.1 0.32 Invention 2 Mn → Si 18 0.014 1.1 0.32 Invention 3 Mn → Si 18 0.014 1.1 0.32 Comparative Material 1 Mn → Si 32 0.014 1.1 0.32 Comparative Material 2 Mn → Si 58 0.014 1.1 0.32 Inventive Steel 2 Invention 4 Mn → Si 16 0.05 1.0 0.35 Invention Steel 3 Invention 5 Mn → Si 15 0.015 1.0 0.35 Inventive Steel 4 Invention 6 Mn → Si 15 0.02 1.0 0.35 Inventive Steel 5 Invention 7 Mn → Si 12 0.05 1.2 0.30 Inventive Steel 6 Invention Material 8 Mn → Si 17 0.02 1.2 0.30 Inventive Steel 7 Invention 9 Mn → Si 18 0.015 1.1 0.32 Inventive Steel 8 Invention 10 Mn → Si 14 0.018 1.1 0.32 Inventive Steel 9 Invention 11 Mn → Si 19 0.02 1.1 0.32 Inventive Steel 10 Invention Material12 Mn → Si 22 0.025 1.0 0.35 Inventive Steel 11 Invention Material 13 Mn → Si 20 0.019 1.0 0.35 Conventional steel (1-11) is not specifically described manufacturing conditions

division Heating temperature (℃) Sedimentary atmosphere (ℓ / min) Heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount / accumulated loading amount of recrystallization zone (%) Cooling rate (℃ / sec) Base nitrogen (ppm) Invention 1 1220 350 160 1030 830 55/75 15 105 Invention 2 1190 610 120 1020 830 55/75 15 115 Invention 3 1150 780 100 1020 830 55/75 15 120 Comparative Material 1 1050 220 50 1020 840 55/75 6 48 Comparative Material 2 1300 950 180 1020 840 55/75 35 420 Invention 4 1180 780 110 1010 830 55/75 16 275 Invention 5 1200 600 100 1040 850 55/75 17 112 Invention 6 1170 620 130 1030 840 55/75 17 80 Invention 7 1190 780 100 1020 830 55/75 16 300 Invention Material 8 1200 620 110 1030 830 55/75 16 100 Invention 9 1150 750 160 1040 830 60/70 16 115 Invention 10 1180 630 110 1040 850 60/70 15 120 Invention 11 1200 520 100 1050 840 60/70 18 90 Invention Material12 1210 550 120 1040 840 60/70 17 100 Invention Material 13 1230 680 110 1030 840 60/70 18 132 Conventional Steel 11 1200 - - Ar ₃ or higher 960 Cooling - The cooling of the invention material controls the cooling rate to 400 ° C., which is the temperature after the bainite transformation is completed, and then air-cooled thereafter. The conventional steel (1-11) is a hot rolled material manufactured without being immersed in a conventional steel. (1-10) does not specify the hot rolling conditions in detail.

Alloying element ratio after nitriding treatment to show the effect of the present invention Ti / N Zr / N N / B Al / N V / N (Ti + Zr + 2Al + 4B + V) / N Invention 1 1.3 1.0 15.0 3.8 1.0 11.1 Invention 2 1.2 0.9 16.4 3.5 0.9 10.1 Invention 3 1.2 0.8 17.1 3.3 0.8 9.7 Comparative Material 1 2.9 2.1 6.9 8.3 2.1 24.3 Comparative Material 2 0.3 0.2 60 1.0 0.2 2.8 Invention 4 1.8 0.7 28.0 2.5 0.4 8.1 Invention 5 1.4 0.9 36.7 5.5 1.8 14.8 Invention 6 2.5 1.3 16.0 2.5 6.3 15.3 Invention 7 1.7 0.7 20.0 3.0 1.7 10.2 Invention Material 8 2.0 1.0 10.0 2.5 9.0 17.4 Invention 9 1.3 0.9 14.4 3.5 1.7 11.1 Invention 10 1.5 0.4 12.0 5.0 0.8 13.1 Invention 11 2.2 1.1 22.5 2.8 2.2 11.3 Invention Material12 2.5 0.5 16.7 4.5 2.0 14.2 Invention Material 13 1.4 0.4 12.0 3.6 - 9.3 Conventional Steel 1 4.1 4.1 13.8 0.6 - 5.7 Conventional Steel 2 2.5 2.5 96.0 0.8 - 4.0 Conventional Steel 3 0.8 0.8 105.8 0.4 - 1.5 Conventional Steel 4 4.1 4.1 4.0 0.8 8.8 15.5 Conventional Steel 5 6.5 6.5 4.0 1.1 18.5 28.1 Conventional Steel 6 3.2 3.2 2.6 0.4 16.1 21.6 Conventional Steel 7 1.0 1.0 9.9 2.5 - 6.5 Conventional Steel 8 1.2 1.2 14.3 0.4 - 2.2 Conventional Steel 9 0.8 0.8 9.1 2.1 3.9 9.2 Conventional Steel 10 0.6 0.6 9.5 3.2 1.5 8.9 Conventional Steel 11 5.5 5.5 12.7 3.4 7.8 20.3

division TiN precipitate properties Base material texture characteristics Base material mechanical properties Count (pcs / mm ² ) Average size (㎛) Thickness (㎛) AGS FGS Ferrite Percentage (%) Thickness (mm) Yield strength (MPa) Tensile Strength (MPa) Elongation (%) -40 ℃ impact toughness (J) Invention 1 2.3 X ^{10 8} 0.016 0.26 17 6 92 20 554 773 35 364 Invention 2 3.1 X ^{10 8} 0.017 0.26 15 5 94 20 595 781 36 355 Invention 3 2.5 X ^{10 8} 0.012 0.24 13 4 93 20 594 780 36 358 Comparative Material 1 4.3X10 ⁶ 0.154 1.4 38 27 70 20 397 584 28 212 Comparative Material 2 5.4 X 10 ⁶ 0.155 1.5 34 23 75 20 394 580 29 189 Invention 4 3.2 X ^{10 8} 0.025 0.35 15 6 93 25 594 784 35 358 Invention 5 2.6 X ^{10 8} 0.013 0.32 14 6 92 25 594 785 35 349 Invention 6 3.3 X ^{10 8} 0.026 0.42 15 6 94 25 591 783 35 230 Invention 7 4.6 X ^{10 8} 0.024 0.45 16 5 93 30 592 784 35 346 Invention Material 8 4.3X10 ⁸ 0.014 0.35 15 6 92 30 595 782 36 352 Invention 9 5.6 X ^{10 8} 0.028 0.36 15 6 91 30 592 786 36 348 Invention 10 5.2 X ^{10 8} 0.021 0.35 15 8 92 30 594 786 35 358 Invention 11 3.7 X ^{10 8} 0.029 0.29 14 7 94 35 591 796 36 362 Invention Material12 3.2 X ^{10 8} 0.025 0.25 16 8 93 35 593 782 35 347 Invention Material 13 3.2 X ^{10 8} 0.024 0.34 15 6 87 35 585 768 36 362 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-TiN 2.80 × 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 0.2μm or less 11.1 × 10 ³ 50 401 514 18.3

As shown in Table 5, the number of precipitates (TiN precipitates) of the hot rolled material produced by the present invention has a range of 1.0 × ^{10 8} / mm ² or more, whereas the precipitates of conventional materials are 4.07X10 ⁵ / mm ² The following range was shown. It was found in the invention that a precipitate of ZrN of 50-100 nm is present. In the present invention, the steel of the present invention is composed of not only fine grain size of ferrite but also a high fraction of ferrite phase.

division Austenitic grain size of welding heat affected zone (㎛) Microstructure with welding heat effect of 100kJ / cm heat input Mechanical Properties of Weldment 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 = 180 seconds Δt _800-500 = 120 seconds Δt _800-500 = 180 seconds Yield strength (kg / ㎡) Tensile Strength (kg / ㎡) Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Invention 1 23 33 56 73 16 370 -74 328 -67 294 -62 Invention 2 22 34 55 76 15 383 -76 343 -69 301 -63 Invention 3 23 32 56 74 17 365 -72 333 -67 298 -63 Comparative Material 1 54 84 182 36 32 126 -43 48 -34 26 -27 Comparative Material 2 65 91 198 37 35 104 -40 39 -32 18 -26 Invention 4 25 37 65 75 18 353 -71 323 -68 287 -64 Invention 5 26 40 57 74 16 362 -71 333 -67 296 -61 Invention 6 25 31 53 76 17 386 -73 352 -69 305 -62 Invention 7 24 34 55 74 18 367 -71 336 -67 293 -63 Invention Material 8 27 36 51 73 14 364 -71 334 -67 294 -61 Invention 9 24 36 52 74 17 367 -72 335 -67 285 -62 Invention 10 22 35 55 73 18 385 -72 340 -66 294 -61 Invention 11 26 34 63 74 16 358 -71 324 -68 285 -63 Invention Material12 27 38 63 74 18 355 -71 323 -67 284 -62 Invention Material 13 24 32 55 75 16 367 -72 334 -68 285 -63 Conventional Steel 1 187 -51 Conventional Steel 2 156 -48 Conventional Steel 3 148 -50 Conventional Steel 4 230 93 143 -48 132 (0 ℃) Conventional Steel 5 180 87 132 -45 129 (0 ℃) Conventional Steel 6 250 47 153 -43 60 (0 degrees Celsius) Conventional Steel 7 141 -54 -61 Conventional Steel 8 156 -59 -48 Conventional Steel 9 145 -54 -42 Conventional Steel 10 138 -57 -45 Conventional Steel 11 141 -43 219 (0 ℃)

Table 6 shows the properties of the weld heat affected zone of the present invention steel and conventional steel. The austenitic grain size at the maximum heating temperature of 1400 ° C. such as the welding heat affected zone shows that the present invention has a range of 52-64 μm, while the conventional material has a very coarse range of about 180 μm or more. Can be. 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 weld heat affected zone by further distributing TiN precipitates that are stable even at high temperatures through the immersion treatment on low-nitrogen steel slabs. The base steel also has a useful effect of providing an improved welded structural steel by forming a ferrite composite.

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: 0.008-0.030%, B: 0.0003-0.01 Making a low nitrogen steel slab satisfying%, W: 0.001-0.2%, P: 0.03% or less, S: 0.003-0.05%, O: 0.01% or less and composed of the remaining Fe and other impurities, Heating the slab at a temperature of 1100-1250 ° C. for 60-180 minutes to make N of 0.008-0.03%, while N and Ti, Zr, B, and Al are immersed to satisfy the following relationship, 1.2≤Ti / N≤2.5, 0.3≤Zr / N≤2.0, 10≤N / B≤40, 2.5≤Al / N≤7, 6.8≤ (Ti + Zr + 2Al + 4B) / N≤17, The heated slab is hot rolled in an austenite recrystallization zone with a reduction ratio of 40% or more, and then cooled by a rate of 5-20 ° C / sec to a bainite transformation end temperature ± 10 ° C. A method for producing a high strength welded steel having a precipitate of TiN and ZrN. 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 + Zr + 2Al + 4B + V) / N. A method for manufacturing a high strength welded structural steel having precipitates of TiN and ZrN by immersion treatment characterized by satisfying ≤19. The steel material according to claim 1 or 2, wherein the steel comprises Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%. High strength welding with precipitates of TiN and ZrN by immersion treatment characterized by containing one or two or more selected and one or two selected from the group Ca: 0.0005-0.005% and REM: 0.005-0.05% Method of manufacturing structural steels. The method of claim 1, wherein the steel slab is added to the molten steel deoxidation element having a greater deoxidizing power than Ti before deoxidation of dissolved oxygen of molten steel to 30ppm or less, and within 10 minutes so that the content of Ti is 0.005-0.2%. A method of manufacturing a high strength welded structural steel having precipitates of TiN and ZrN by immersion treatment, characterized in that it is made by continuous casting after addition. The steel material according to any one of claims 1, 2, and 4, wherein the steel is composed of a composite structure of 30-80% bainite and ferrite of 20 micrometers or less, and a ZrN precipitate having a size of 0.01-. A method for producing a high strength welded structural steel having a precipitate of TiN and ZrN by immersion treatment, characterized in that TiN precipitates of 0.1 μm are distributed at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less.