KR100481365B1 - Method of manufacturing steel plate to be precipitating TiN and TiO 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 distribute TiN and TiO oxides more finely and to weld from medium to high heat input. The difference in the toughness of the base material and the heat affected zone is minimized even in the heat input range, thereby providing a method of manufacturing a welded structural steel having excellent toughness of the heat affected zone.

The present invention for achieving the above object, the primary deoxidation element is added to molten steel to adjust the dissolved oxygen amount to 60 to 120ppm after the second deoxidation with Ti deoxidizer and the third deoxidation to add a deoxidation element having a higher deoxidizing power than Ti. After adjusting the dissolved oxygen amount to 60ppm or less, after adding Ti and refining the amount of Ti to 0.005-0.2%, 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: 0.03% or less, O: 0.001-0.006%, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10 Casting into a steel slab composed of Fe and other impurities,

The steel slab was heated in the range of 1100-1250 ° C. for 60-180 minutes and then hot-rolled in the austenite recrystallization zone with a rolling ratio of 40% or more, and then cooled at a rate of 1 ° C./min or more until the end of ferrite transformation temperature ± 10 ° C. Technical aspect of the present invention relates to a method for manufacturing a welded structural steel having a fine TiO oxide and TiN precipitates.

Description

Method for manufacturing steel plate to be precipitating TiN and TiO for welded structures}

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 distributing TiN precipitates and TiO oxides.

Recently, the steel used in accordance with the trend of high-rise building, structure has been replaced by thick steel.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

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

Until now, many techniques have been known to improve the toughness of weld heat affected zones during high heat input welding using TiN precipitates and Al-based or MgO oxides. No significant improvements have been made yet. 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.

The present invention is a welded structural steel having excellent toughness of the heat-affected portion while minimizing the difference in the toughness of the base material and the heat-affected portion even in the weld heat input range ranging from medium heat to superheated heat by finely distributing TiN and TiO oxides stable at high temperature. To provide a method for the preparation, the purpose is.

In the method of manufacturing the welded structural steel of the present invention for achieving the above object, the primary deoxidation element is added to molten steel to adjust the dissolved oxygen content of the molten steel to 60 to 120 ppm, followed by secondary deoxidation with Ti deoxidizer and higher deoxidizing power than Ti. After adjusting the dissolved oxygen to 60ppm or less by tertiary deoxidation with a large deoxidation element, after adding Ti and refining the amount of Ti to 0.005-0.2%, 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: 0.03% O: 0.001-0.03%, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti Casting into a steel slab that satisfies / O≤10 and is composed of the remaining Fe and other impurities,

The steel slab was heated in the range of 1100-1250 ° C. for 60-180 minutes and then hot-rolled in the austenite recrystallization zone with a rolling ratio of 40% or more, and then cooled at a rate of 1 ° C./min or more until the end of ferrite transformation temperature ± 10 ° C. It is configured to include.

Hereinafter, the present invention will be described in detail.

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

The present inventors have found that the toughness of the weld heat affected zone changes when the grain size of the austenite is fine (about 80 μm or less) in the course of research for improving the toughness of the weld heat affected zone.

Based on these studies, in the present invention,

[1] with fine distribution of TiN precipitates and TiO oxides,

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

[3] Effective use of TiO oxides and BN, AlN, and VN precipitates increases ferrite fraction in weld heat affected zones, and in particular, promotes the transformation of polygonal or needle-like ferrites in the austenite to enhance toughness improvement. These [1] [2] [3] are demonstrated in more detail.

[1] microdispersion of TiN precipitates and TiO oxides

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. That is, when the ratio of Ti and N (Ti / N) is in the range of 1.2 to 2.5, the amount of solid solution Ti is extremely reduced and the high temperature stability of the TiN precipitate is increased, so that the fine TiN precipitate of 0.01-0.1 μm or less is 0.5 μm or less. A result of distribution of 1.0 × 10 ⁷ pieces / mm 2 or more at intervals was obtained. 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 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 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 deoxidizes in steel, improves weldability, hot workability and strength, and forms a solid solution in the matrix to strengthen the matrix to secure the strength and toughness. To this end, Mn is preferably added at least 0.4%. Do. However, when the Mn content is more than 2.0%, it has a detrimental effect on the toughness of the weld heat affected zone due to tissue heterogeneity due to manganese segregation rather than a solid solution strengthening effect. In addition, when the steel solidifies, macro segregation and micro segregation occur depending on the segregation mechanism, which promotes the formation of a central segregation zone in the center of rolling, thereby causing a low temperature transformation structure in the center of the base metal. In particular, manganese is an element that precipitates in the form of MnS around the Ti-based oxide and affects the formation of acicular and polygonal ferrites effective for improving the toughness of the weld heat affected zone.

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

Al is 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 an indispensable element in the present invention because it forms fine Ti oxide by reacting with O in steelmaking process and also forms fine TiN precipitate by combining with N in the solidification process and slab heating process. In order to obtain such a fine TiO and TiN precipitation effect, it is preferable to add more than 0.005% of Ti, but if it exceeds 0.2%, coarse Ti oxide is formed in molten steel, which does not inhibit the growth of the weld heat-affected part of the austenite grain and improves the toughness. It is not preferred because it does not promote favorable acicular ferrite transformation.

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; Particularly, the content of TiN and AlN precipitates and the spacing of precipitates, the distribution of precipitates, the frequency of complex precipitation with oxides, and the high temperature stability of the precipitates themselves are significantly affected. However, when the nitrogen content exceeds 0.03%, the effect is saturated, and toughness decreases due to an increase in the amount of solid solution nitrogen distributed in the weld heat affected zone, and it may be incorporated into the weld metal due to dilution during welding, resulting in a decrease in the toughness of the weld metal. Can be.

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

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

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

P is preferably as low as possible because it is an impurity element that promotes central segregation during rolling and hot cracking during welding. In order to improve the toughness of the base metal, the toughness of the weld heat affected zone, and to reduce the center segregation, it is recommended to manage it to 0.03% or less.

Since S forms a low melting point compound such as FeS when present in a large amount, it is preferable to manage S as low as possible. In order to reduce the base material toughness, weld heat affected zone toughness and central segregation, it is recommended that the S content be 0.03% or less. In particular, sulfur is precipitated in the form of MnS around Ti-based oxides, which affects the formation of needle-shaped and polygonal ferrites, which are effective for improving the toughness of welding heat affected zones. It is desirable to limit it to 0.03% or less.

The content of oxygen (O) is preferably limited to 0.001-0.006%.

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 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. Affects bad toughness. 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.

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.

Moreover, it is preferable to make ratio of V / N into 0.3-9. In the present invention, when the V / N ratio is less than 0.3, it is difficult to secure an appropriate number and size of VN precipitates deposited and distributed 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.

[Method of manufacturing welded structural steel]

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

According to the study of the present invention, the continuous casting speed is less than 1.2 m / min at low speed, more preferably about 0.9 to 1.2 m / min. The reason is that when the casting speed is less than 0.9m / min, it is advantageous for cast surface cracks, but productivity is lowered. When it is faster than 1.2m / min, cast surface cracks are more likely to occur.

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

Hot rolling process

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

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

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

Precipitate

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

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.

Casting of the steel in the present invention can be produced by slab by continuous casting or mold casting. In this case, if the cooling rate is fast, it is advantageous to finely disperse the precipitates, and thus, continuous casting having a high cooling rate is preferable. For the same reason, the slab is advantageously thinner. In 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 known 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 dissolving them in a converter, and the composition ratio between the alloying elements for each steel type to show the effect of the present invention is shown in Table 2. Table 3 shows the solidification rate of slab, slab heating temperature, heating time, rolling start temperature and end temperature, rolling reduction, and rolling rate of 25 ~ 40mm thickness in rolling process. At this time, the rolling reduction ratio of all steel grades was made 75% or more.

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 0.12 0.13 1.54 0.006 0.005 0.04 0.014 7 115 0.005 - - - - - 0.01 - - 19 Inventive Steel 2 0.08 0.12 1.50 0.006 0.005 0.07 0.05 10 275 0.002 - 0.2 - - - 0.01 - - 50 Invention Steel 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015 3 112 0.003 0.1 - - - - 0.02 - - 21 Inventive Steel 4 0.10 0.12 1.48 0.006 0.005 0.02 0.02 5 80 0.001 - - - - - 0.05 - - 23 Inventive Steel 5 0.09 0.15 1.52 0.006 0.004 0.09 0.05 15 300 0.002 0.1 - 0.1 - - 0.05 - - 50 Inventive Steel 6 0.10 0.14 1.50 0.007 0.005 0.025 0.02 10 100 0.004 - - - 0.1 - 0.09 - - 22 Inventive Steel 7 0.13 0.14 1.48 0.007 0.005 0.04 0.015 8 115 0.15 0.1 - - - - 0.02 - - 21 Inventive Steel 8 0.11 0.15 1.52 0.007 0.005 0.06 0.018 10 120 0.001 - - - - 0.015 0.01 - - 20 Inventive Steel 9 0.13 0.21 1.50 0.007 0.005 0.025 0.02 4 90 0.002 - - 0.1 - - 0.02 0.001 - 21 Inventive Steel 10 0.08 0.16 1.45 0.008 0.006 0.045 0.025 6 100 0.05 - 0.3 - - 0.01 0.02 - 0.01 27 Inventive Steel 11 0.12 0.24 1.48 0.007 0.003 0.047 0.019 11 132 0.01 - 0.1 - - - - - - 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 58 0.009 0.06 38 0.014 15 21 1.0 0.32 Inventive Steel 4 Invention 6 Mn → Si 54 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 55 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

Alloy element composition ratio for showing the effect of the present invention Ti / O Ti / N N / B Al / N V / N (Ti + 2Al + 4B + V) / N Invention 1 7.7 1.2 16.4 3.5 0.9 9.3 Invention 2 7.4 1.2 16.4 3.5 0.9 9.3 Invention 3 7.7 1.2 16.4 3.5 0.9 9.3 Comparative Material 1 17.5 1.2 16.4 3.5 0.9 9.3 Comparative Material 2 1.1 1.2 16.4 3.5 0.9 9.3 Invention 4 10.0 1.8 27.5 2.5 0.4 7.4 Invention 5 7.5 1.3 37.3 5.4 1.8 14.0 Invention 6 8.7 2.5 16.0 2.5 6.3 14.0 Invention 7 10.0 1.7 20.0 3.0 1.7 9.5 Invention Material 8 9.1 2.0 10.0 2.5 9.0 16.4 Invention 9 8.8 1.3 14.4 3.5 1.7 10.3 Invention 10 9.0 1.5 12.0 5.0 0.8 12.7 Invention 11 9.5 2.2 22.5 2.8 2.2 10.2 Invention Material12 9.3 2.5 16.7 4.5 2.0 13.7 Invention Material 13 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

Steel grade used division Heating temperature (℃) Heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount / accumulated loading amount at recrystallization area (%) Cooling rate (℃ / min) Cooling end temperature (℃) Invention 2 Inventive Example 1 1150 170 1030 780 65/80 7 550 Inventive Example 2 1200 130 1040 790 65/80 7 550 Inventive Example 3 1240 90 1040 780 65/80 7 550 Comparative Example 1 1050 60 1040 780 65/80 7 550 Comparative Example 2 1300 250 1035 780 65/80 7 550 Invention 1 Inventive Example 4 1200 130 1020 790 65/80 6 550 Invention 3 Inventive Example 5 1200 130 1040 790 65/80 6 550 Comparative Material 1 Comparative Example 3 1210 120 1030 780 65/80 0.1 Room temperature Comparative Material 2 Comparative Example 4 1210 120 1030 790 65/80 19 Room temperature Invention 4 Inventive Example 6 1180 150 1020 790 60/80 7 550 Invention 5 Inventive Example 7 1190 140 1010 800 60/80 8 550 Invention 6 Inventive Example 8 1220 110 1010 810 60/75 7 550 Invention 7 Inventive Example 9 1220 110 1020 800 60/75 10 550 Invention Material 8 Inventive Example 10 1210 120 1010 790 60/75 10 550 Invention 9 Inventive Example 11 1220 110 1000 780 55/70 10 550 Invention 10 Inventive Example 12 1210 120 1010 790 55/70 9 550 Invention 11 Inventive Example 13 1230 100 1000 800 55/70 8 550 Invention Material12 Inventive Example 14 1220 110 1020 780 55/70 10 550 Invention Material 13 Inventive Example 15 1210 130 1020 780 65/75 10 550 Conventional Steel 11 1200 - Ar ₃ or higher 960 80 Cooling Manufacturing conditions of conventional steel (1-10) are not specifically presented

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 at -40 ° C (J) Inventive Example 1 2.5 X ^{10 8} 0.015 0.24 2.3 X 10 ² 1.8 25 394 553 38 11 82 359 Inventive Example 2 3.3 X ^{10 8} 0.013 0.23 2.2 X 10 ² 1.4 25 395 551 39 9 83 363 Inventive Example 3 2.7 X ^{10 8} 0.011 0.22 2.4 X 10 ² 1.2 25 396 550 39 10 83 358 Comparative Example 1 2.4 X 10 ⁶ 0.174 1.6 2.5X10 ¹ 2.2 25 393 554 26 16 70 206 Comparative Example 2 3.4 X 10 ⁶ 0.165 1.8 2.2 X 10 ¹ 2.3 25 792 860 17 17 21 45 Inventive Example 4 3.5 X ^{10 8} 0.015 0.30 3.3 X 10 ² 1.2 30 396 558 38 11 83 352 Inventive Example 5 2.8X10 ⁸ 0.013 0.31 2.7 X 10 ² 1.2 30 396 562 38 10 83 357 Comparative Example 3 1.3 X 10 ⁶ 0.182 1.2 1.4 X 10 ¹ 3.4 30 384 564 30 18 73 220 Comparative Example 4 4.3X10 ⁶ 0.177 1.4 1.1 X 10 ¹ 3.7 30 392 582 29 17 74 208 Inventive Example 6 3.5 X ^{10 8} 0.016 0.33 1.9 X 10 ² 1.3 30 390 563 38 10 82 365 Inventive Example 7 4.8 X ^{10 8} 0.014 0.30 3.3 X 10 ² 1.2 35 390 564 39 10 85 362 Inventive Example 8 4.4 X ^{10 8} 0.014 0.39 2.2 X 10 ² 1.6 35 392 542 36 11 82 367 Inventive Example 9 5.7 X ^{10 8} 0.018 0.27 2.5 X 10 ² 1.5 35 391 536 37 10 84 363 Inventive Example 10 5.4 X ^{10 8} 0.011 0.23 2.3 X 10 ² 1.5 35 394 566 36 10 83 378 Inventive Example 11 3.9 X ^{10 8} 0.019 0.23 2.7 X 10 ² 1.7 40 390 566 37 12 83 365 Inventive Example 12 3.4 X ^{10 8} 0.015 0.30 1.7 X 10 ² 1.5 40 396 542 38 11 85 358 Inventive Example 13 3.6X10 ⁸ 0.012 0.33 1.9 X 10 ² 1.5 40 406 564 38 12 82 349 Inventive Example 14 3.8X10 ⁸ 0.012 0.27 2.3 X 10 ² 1.6 40 387 550 37 10 83 354 Inventive Example 15 4.4 X ^{10 8} 0.018 0.25 2.0X10 ² 1.7 30 389 549 39 9 86 373 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 Conventional Steel 5 Precipitates of MgO-TiN 4.07 × 10 ⁶ pcs / mm ² 40 494 622 Conventional Steel 6 Precipitates of MgO-TiN 2.80 × 10 ⁶ pcs / mm ² 50 812 912 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 50 401 514

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 2.5 × ^{10 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. Meanwhile, it can be seen that the ferrite grain size (FGS) of the inventive steel in the base steel structure of the inventive steel is very fine compared to that of the comparative material in the range of about 9-12 μm. 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 54 75 15 373 -74 337 -67 295 -63 Inventive Example 2 22 35 53 78 13 385 -76 352 -69 307 -64 Inventive Example 3 19 35 54 75 13 368 -72 333 -67 296 -63 Comparative Example 1 54 86 182 38 24 123 -43 43 -34 28 -28 Comparative Example 2 65 92 198 36 26 101 -40 30 -32 17 -25 Inventive Example 4 22 38 61 76 14 354 -71 330 -68 289 -65 Inventive Example 5 24 41 55 78 15 367 -71 336 -67 299 -62 Comparative Example 3 56 80 178 40 26 103 -39 56 -32 24 -24 Comparative Example 4 63 88 184 39 28 66 -28 39 -30 10 -21 Inventive Example 6 19 32 50 75 14 384 -73 357 -69 313 -63 Inventive Example 7 21 35 51 77 14 366 -71 341 -67 298 -63 Inventive Example 8 27 37 52 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 52 75 14 384 -72 346 -66 294 -63 Inventive Example 11 26 35 61 75 14 356 -71 329 -68 286 -68 Inventive Example 12 27 39 59 74 15 354 -71 323 -67 278 -62 Inventive Example 13 23 38 58 74 14 355 -71 323 -67 255 -62 Inventive Example 14 25 35 52 70 15 344 -71 328 -67 249 -63 Inventive Example 15 23 36 49 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 ℃)

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 uses a TiO oxide together with TiN precipitates to develop a welded structural steel material having excellent base material properties and at the same time securing excellent weld heat affected zone properties. By controlling the grains and promoting the needle ferrite transformation in the grains it can provide a structural structural steel for welding that can attain a good heat input weld heat affected zone toughness at the same time.

Claims (6)

After adjusting the dissolved oxygen amount of molten steel to 60 ~ 120ppm by adding deoxidation element to molten steel, and adjusting the dissolved oxygen amount to 60ppm or less by tertiary deoxidation with Ti deoxidizer and adding deoxidation element with more deoxidizing power than Ti. Ti was refined to add 0.005-0.2% Ti, followed by 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: 0.03% or less, O: 0.001-0.006%, 1.2≤Ti / N≤2.5, 10≤ Steel slab that satisfies N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10 and is composed of the remaining Fe and other impurities Molding step, The steel slab was heated in the range of 1100-1250 ° C. for 60-180 minutes and then hot-rolled in the austenite recrystallization zone with a rolling ratio of 40% or more, and then cooled at a rate of 1 ° C./min or more until the end of ferrite transformation temperature ± 10 ° C. Method for producing a welded structural steel having a fine TiO oxide and TiN precipitates comprising a. 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. Method for producing a welded structural steel having a fine TiO oxide and precipitates of TiN. 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%. Manufacture of welded structural steel having fine TiO oxide and TiN precipitates, 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 to 0.05% Way. The method of claim 1, wherein the casting is performed by continuously casting the molten steel at a speed of 0.9 to 1.2 m / min of the fine TiO oxide and TiN in the secondary cooling zone at a specific amount of 0.3 ~ 0.35 l / kg Method for producing welded steel having precipitates. The welding structure having a fine TiO oxide and TiN precipitate according to claim 1, wherein the tertiary deoxidation is carried out within 10 minutes after the secondary deoxidation with Ti so that the final content of Al is 0.005-0.1%. Method of manufacturing steel. 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 , 0.5-2.0㎛ size of TiO oxide 1 × 10 ² -1 × 10 ³ / mm ^{2 The} method of manufacturing a welded structural steel having a fine TiO oxide and TiN precipitates, characterized in that distributed.