KR100470053B1 - Steel plate to be precipitating TiN and complex oxide of Mg-Ti, method for manufacturing the same, welding fabric using the same - 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 the toughness of the base material and the heat affected zone using TiN precipitates and Ti-Mg composite oxides. It is to provide a welded structural steel material having excellent toughness of the heat-affected zone and a manufacturing method thereof, and a welded structure using the heat-affected zone characteristic of the structural steel while the difference is minimized.

The present invention for achieving the above object, in the weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1%, N: 0.008-0.030%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 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, 0.2≤Mg / O≤3, 3 ≤ (Ti + Mg) / O ≤ 12, is composed of the remaining Fe and other impurities, and has a TiN precipitate and Mg-Ti composite oxide consisting of a composite structure of ferrite and pearlite of 20 micrometers or less Technical aspects of welded structural steels shall be taken as their technical gist.

Description

Steel plate to be precipitating TiN and complex oxide of Mg-Ti, method for manufacturing the same, welding fabric using the same}

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 welded structural steel material and a method of manufacturing the same, by using fine TiN precipitates and Ti-Mg composite oxides to improve the toughness of the weld heat affected zone.

Recently, the steel used in accordance with the trend of high-rise building, structure has been replaced by thick steel.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 steel materials, a large-pass heat submerged welding method and an electro-welding method capable of one-pass welding are widely used. In addition, in the field of shipbuilding and bridges, even when welding a steel plate having a plate thickness of 25 mm or more, the above-described high heat input welding method capable of one-pass welding is applied.

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

When the heat input is applied to the steel, the heat affected zone formed during welding, particularly the heat affected zone near the fusion boundary, is heated to a temperature close to the melting point by the amount of heat input. Accordingly, since the grains of the weld heat affected zone grow and coarse, and microstructures that are vulnerable to toughness such as upper bainite and martensite are formed during the cooling process, the weld heat affected zone is the site where the toughness of the weld deteriorates most.

Therefore, in order to secure the stability of the welded structure, it is necessary to suppress the growth of the austenite grains in the weld heat affected zone and to keep it fine. As a means to solve this problem, there is disclosed a technique for delaying grain growth of the weld heat affected zone during welding by appropriately distributing an oxide or Ti-based carbonitride, which is stable at a high temperature, to steel materials. For example, Japanese Patent Application Laid-Open No. 11-140582, No. 10-298708, No. 10-298706, No. 9-194990, No. 9-324238, No. 8-60292, (S) 60-245768, (Pyeong) 5-186848, (S) 58-31065, (S) 61-79745, Journal of the Japan Welding Society, Vol. 52, No. 2, 49 and Japanese Patent Laid-Open And 64-15320.

Japanese Patent Laid-Open No. 11-140582 is a representative technique using TiN precipitates, and has a toughness of about 200J at 0 ° C when 100 J / cm of heat input (maximum heating temperature of 1400 ° C) is applied. Is about 300J). In this prior art, Ti / N is substantially managed at 4-12 so that TiN precipitates of 0.05 µm or less are 5.8 × 10 ³ pieces / mm 2 to 8.1 × 10 ⁴ pieces / mm 2, and TiN precipitates of 0.03 to 0.2 μm are 3.9 ×. The toughness of the welded portion is secured by making the ferrite fine by depositing 10 ³ pieces / mm 2 to 6.2 × 10 ⁴ pieces / mm 2. However, according to this prior art, when the 100 kJ / cm high heat input welding is applied, the toughness of the base material and the heat affected zone is generally low (the base material: 320J, the heat affected zone: 220J at the highest impact toughness of 0 ° C), and also the base material. Since the toughness difference between the and the heat affected zone is about 100J, there is a limit in securing the reliability of the steel structure due to superheated welding of the thickened steel. In addition, as a method for securing the desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or higher and quenched, followed by a reheating process for hot rolling. Becomes

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 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.

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- or MgO-TiN composite precipitates. No significant case has yet been announced for improving the toughness of the affected area. 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.

In the present invention, using the TiN precipitate having excellent high temperature stability and the composite oxide of Ti-Mg, the toughness of the base material and the heat affected zone is minimized while the toughness of the heat affected zone is excellent, and the manufacturing method of the welded structural steel and its structural steel, and It is an object of the present invention to provide a welded structure using heat affected zone characteristics.

Welded structural steel materials 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: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 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, 0.2≤Mg / O It satisfies ≦ 8, 3 ≦ (Ti + Mg) / O ≦ 12, is composed of the remaining Fe and other impurities, and the microstructure is composed of a composite structure of ferrite and pearlite of 20 μm or less.

In addition, the manufacturing method of the welded structural steel material of this invention is 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% by weight. , 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: 0.002-0.03%, Mg: 0.001-0.005%, 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, 0.2≤Mg / O≤ 3, 3≤ (Ti + Mg) / O≤12, and the slab composed of the remaining Fe and other impurities is heated in the range of 1000-1250 ° C for 60-180 minutes, followed by a rolling ratio of 40% or more in the austenite recrystallization zone. After hot rolling, the ferrite transformation end temperature includes a cooling at a rate of 1 ° C / min or more.

In addition, the welded structure of the present invention, 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: 0.002-0.03%, Mg: 0.001-0.005%, 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, 0.2≤Mg / O≤8, 3≤ (Ti + Mg) / O ≦ 12, composed of the remaining Fe and other impurities, 1.0 × 10 ⁷ TiN precipitates of 0.01-0.1 μm in a composite structure of ferrite and pearlite of 20 μm or less at intervals of 0.5 μm / Weld is applied to steel materials (base metal) with Ti-Mg composite oxide of 2.0 ~ 5.0㎛ distributed in mm2 or more and 1 × 10 ² -1 × 10 ³ pcs / mm ² . Austenite is formed and then quenched so that the microstructure of the weld heat affected zone is 20 micrometers or less of ferrite with a percentage of 70% or more.

Hereinafter, the present invention will be described in detail.

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

In the course of research to improve the toughness of the weld heat affected zone, when the TiN precipitate and Mg-Ti composite oxide are used together, the crystal grain size of the austenite of the weld heat affected zone is about 80 μm or less due to the main action of TiN. It has been found that the micronization and the main action of the composite oxide of Mg-Ti promote the nucleation of needle-like ferrite from fine guustenite, thereby significantly improving the toughness of the weld heat affected zone.

In the present invention starting from this point of view,

[1] using TiN precipitate and Ti-Mg composite oxide together,

[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 in the heat affected zone is refined to 80 µm or less. Also,

[3] Improved toughness by increasing Ti-Mg composite oxide, BN, and AlN precipitates to increase the ferrite formation rate from the former austenite in the weld heat affected zone, and in particular, promote the transformation of polygonal or needle-like ferrite in the old austenite. Increase the effect These [1] [2] [3] are demonstrated in more detail.

[1] management of TiN precipitates and Ti-Mg composite 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 present inventors focused on the fact that the TiN precipitates distributed in the base metal are caused by the diffusion of solid solution Ti atoms decomposed by welding heat, and consequently examine the characteristics of TiN precipitates according to the ratio of Ti / N. At low Ti / N ratio, new findings show that the dissolved Ti concentration, the 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. Increasing the nitrogen content at the same Ti content, all of the Ti atoms dissolved are easily combined with the nitrogen atom, and in the high nitrogen environment, the TiN precipitate is more stable at high temperature than in the case where the nitrogen content is low. The degradation was analyzed to be due to the lower solubility product. 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.

Furthermore, in the present invention, the Ti-Mg composite oxide, which is stable at 1400 ° C. or more, together with the fine and uniformly distributed TiN precipitates, has another feature. The Ti-Mg composite oxide is dispersed in the base material and welded during welding. In addition to suppressing the growth of the affected zone of austenite grains, the toughness of the weld heat affected zone can be improved by promoting the needle-like ferrite transformation from the austenite. In the precipitation of Ti-Mg complex oxide, the ratio of Mg, Ti and O has an important influence.

[2] ferrite grain size management

According to the study of the present invention, in order to make the size of the old austenite at 80 μm in the weld heat affected zone, it is important to reduce the size of the ferrite to about 20 μm or less in the base material structure of ferrite + pearlite together with the management of precipitates. It is. At this time, the refinement of the ferrite is obtained by the growth inhibition of the ferrite grains generated in the cooling process by using carbide (Fe ₃ C, VC, WC) with the austenite grain refinement by the steel processing during hot rolling.

[3] microstructure of weld heat affected zone

The facts of the present invention reveal that the toughness of the weld heat affected zone is not only the size of the former austenite grains when the base material is heated above about 1400 ° C., but also the amount of ferrite (more than 70%) precipitated at the old austenite grain boundaries. The size (less than 20㎛) and its shape have an important influence. In the present invention, the Mg / O ratio is 0.2-3, the Ti / O ratio is 4-10, and the (Ti + Mg) / O ratio is 3-12. In addition to the appropriate ratio of Mg, using a precipitate of AlN, BN, etc. to generate a large amount of fine ferrite in the mouth of the austenite, the ferrite is mostly polygonal ferrite and needle-like ferrite toughness of the heat affected zone Greatly improves.

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

[Welding Structural Steels]

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

When the content of carbon (C) is less than 0.03%, securing strength as a structural steel is insufficient. In addition, when C exceeds 0.17%, microstructures susceptible to toughness, such as upper bainite, martensite and degenerate pearlite, are transformed during cooling to lower the low temperature impact toughness of structural steel, and also the hardness of the welded portion. Or increasing the strength resulting in deterioration of toughness and generation of weld cracks.

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

If the content of silicon is less than 0.01%, the deoxidation effect of molten steel is insufficient during steelmaking and the corrosion resistance of steel is reduced. If the content is more than 0.5%, the effect is saturated, It promotes the transformation of martensite and lowers the low temperature impact toughness.

The content of manganese (Mn) is preferably limited to 0.4-2.0%.

Mn is an effective element that improves deoxidation, weldability, hot workability and strength in steel, and precipitates in the form of MnS around Ti-based oxides, which affects the formation of acicular and polygonal ferrites effective for improving the toughness of weld heat affected zones. Element. Such Mn forms a solid solution in the matrix structure to strengthen the matrix and solidify the matrix to secure strength and toughness. However, when the Mn content exceeds 2.0%, tissue heterogeneity due to Mn segregation has a detrimental effect on the toughness of the weld heat affected zone, rather than the effect of strengthening the solid solution. In addition, when the Mn content is added more than 2.0%, macro segregation and micro segregation occurs depending on the segregation mechanism during steel coagulation, which promotes the formation of a central segregation zone in the center part during rolling, thereby creating a low temperature transformation structure in the center of the base material. .

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

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 combines with N to form a fine TiN precipitate that is stable at high temperature as well as to form a Ti-Mg complex oxide. In order to obtain such a fine TiN and Ti-Mg composite precipitation effect, it is preferable to add more than 0.005% of Ti, but when it exceeds 0.2%, coarse TiN precipitates and coarse Ti-Mg composite oxides are formed in molten steel. And the austenite grain growth of the weld heat-affected zone at the time of welding because it is mixed in the base material and is not preferable.

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 suppresses the growth of the old austenite grains at the weld heat affected zone during the high heat input welding and increases 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 0.0020-0.03%.

O is an element that reacts with Ti and Mg in molten steel to form a Ti-Mg composite oxide. The Ti-Mg composite oxide promotes the transformation of acicular ferrite from guustenite in the weld heat affected zone. If the O content is less than 0.0020%, the desired Ti-Mg composite oxide cannot be obtained. If the O content exceeds 0.03%, coarse Ti-Mg composite oxide and other oxides such as FeO are produced, which is not preferable.

The content of magnesium (Mg) is preferably limited to 0.001-0.005%.

Mg is a useful element that combines with Ti and O in molten steel to form Ti-Mg composite oxides to promote needle ferrite transformation in the weld heat affected zone. Mg content of 0.001% or more is required to form fine Ti-Mg composite oxides. It is preferable to add. However, when the Mg content exceeds 0.005%, coarse Mg oxide is formed in molten steel, which adversely affects the mechanical properties of the base metal.

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 Ti atoms that are dissolved in the cooling process during the playing process are combined with the nitrogen atoms to increase 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, not only the number of Ti-Zr oxides required for austenite grain growth inhibition is insufficient, but also the Ti ratio contained in the oxides becomes small, resulting in loss of function as a needle-like ferrite nucleation site. The acicular ferrite phase fraction effective for improving the toughness of the heat affected zone is lowered. When the ratio of Ti / O is over 10, the effect of inhibiting austenite grain restraint on the welding heat-affected zone is saturated, and the proportion of sulfides such as MnS contained in Ti oxide is reduced, thereby reducing the function of ferrite nucleation sites in the mouth. Loss

The Mg / O ratio is preferably limited to 0.2-3.

If the Mg / O ratio is less than 0.2, the number of Ti-Mg complex oxides required for austenite grain growth inhibition is insufficient, and the Mg ratio contained in the oxide becomes small, and thus loses its function as a needle-like ferrite nucleation site. The needle ferrite phase fraction effective for improvement is lowered. In addition, when the Mg / O ratio exceeds 3, the austenite grain growth inhibitory effect is saturated, and the ratio of components such as Mn contained in the oxide becomes smaller, thus losing the function of nucleation sites of the ferrite in the mouth.

The (Ti + Mg) / O ratio is preferably limited to 3-12.

When the (Ti + Mg) / O ratio is less than 3, the number of Ti-Mg complex oxides required for the austenite grain growth inhibition is insufficient, and thus the toughness of the weld heat affected zone is improved by losing its function as a needle ferrite nucleation site in the austenite grain. The effective acicular ferrite phase fraction decreases. In addition, when the (Ti-Mg) / O ratio exceeds 12, the austenite grain growth inhibitory effect is saturated, and the ratio of the Mn component 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 the proper number and size of VN precipitates deposited and distributed at the Ti-Mg oxide + MnS precipitate boundary for improving the toughness of the weld heat affected zone. If the V / N ratio exceeds 9, the weld size is reduced because the VN precipitates deposited on the Ti-Mg oxide + MnS precipitate boundary are coarsened and the number of VN precipitates deposited on the Ti-Mg oxide + MnS complex precipitate decreases. Reduces the ferrite phase fraction effective for the toughness of the heat affected zone.

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 that enhances the strength and toughness of the base material by strengthening the solid solution. In order to obtain such an effect, Ni is preferably contained in an amount of 0.1% or more. It is not preferable because it lowers the toughness and there is a possibility of hot cracking in the weld heat affected zone and the weld metal.

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 the REM exceeds 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.

· Microstructure of steel

In the present invention, the microstructure of the steel is ferrite + pearlite, the size of the ferrite grains is preferably 20㎛ 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 reusable 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, it is preferable to make the distribution of TiN uniform by controlling the space | interval of TiN precipitate to 0.5 micrometer or less.

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

In the present invention, the particle diameter and the critical number of the Ti-Mg oxide are preferably in the range of 0.5-2.0 μm and 1.0 × ^{10 2} -1.0 × 10 ³ per 1 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 oxides is less than 1.0 × ^{10 2} per 1 mm ² , the number of oxides nucleating the needle ferrite in the weld heat affected zone is small and the amount of needle ferrite is low, which adversely affects the toughness and exceeds 1.0 × 10 ³ per 1 mm ² . In this case, since the number of Ti-Mg complex oxides increases, the needle-like ferrites nucleating in the oxides grow and collide with each other to limit the formation of the secondary needle-like ferrites.

[Method of manufacturing welded structural steel]

Refining (Deoxidation, Degassing) Process

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

The present inventors focused on the fact that the amount of dissolved oxygen in the deoxidation process greatly affects the formation behavior of the oxide to uniformly distribute a large amount of Ti into an appropriate amount of Ti-Mg composite oxide and a large amount of fine TiN precipitate. The following facts could be found.

(1) The amount of dissolved oxygen greatly affects the formation behavior of oxides, and in order to generate a large amount of oxides, an appropriate amount of dissolved oxygen exists, which is about 50-200 ppm,

(2) When Mg is first introduced into molten steel having such dissolved oxygen, and then Ti is added, Mg oxide has a significantly higher affinity with oxygen than other oxides, and thus a finer oxide distribution can be obtained. Some form Ti-Mg complex oxides to increase the number of oxides, and these Ti-Mg complex oxides increase the precipitation frequency of precipitates such as TiN, MnS, CuS, VN, and the rest of the added Ti in molten steel. The solid solution forms a fine Ti precipitate upon solidification, which will be described in detail.

In the present invention, before the addition of Mg and Ti, an element having a higher deoxidation power than Ti is added to molten steel to adjust the dissolved oxygen amount to 50-200 ppm. The deoxidizing power of the deoxidizer is as follows.

Cr <Mn <Si <Ti <Al <REM <Zr <Ca ≒ Mg

The deoxidizer with greater affinity with oxygen has a higher rate of binding to oxygen in molten steel and a stronger binding force. Therefore, before adding Mg and Ti, a deoxidizer having a higher deoxidizing power than Ti must be added to secure an appropriate dissolved oxygen (50-200 ppm) to form TiN as a TiN precipitate while forming a target TiO oxide while reducing the progress time of the deoxidation process. Can be. That is, since the deoxidizer with greater deoxidizing power than Ti is pre-injected, the ratio of Ti remaining in solid solution rather than coarse primary oxide in molten steel increases considerably, thus securing fine amounts of Ti oxide and TiN nitride during solidification. It can be.

In the present invention, before adding an element having a higher deoxidizing power than Ti (Al, REM, Zr, Ca, Mg), by deoxidizing by adding elements such as Mn, Si, and then deoxidizer having a higher deoxidizing power than Ti such as Al It is also preferable to inject. This method has the advantage of adjusting the amount of oxygen of molten steel through the separation of coarse primary oxide by adjusting Al and adding Al and adding Al by adding weak Al.

As an example of the deoxidation operation pattern in the present invention, Mn, Si, Al are added in the order of primary refining, followed by Mg in the secondary refining, and finally Ti is added. In the present invention, the final Al content is limited to the range of 0.005-0.1%, because the amount of oxygen in the molten steel is reacted with the oxygen in the molten steel by the addition of Al to be added after the first deoxidation in the above-described content range. Because it is effective to control. As a result of experiment to derive the proper range of Al, when Al is less than 0.005%, the influence on the control of the proper dissolved oxygen amount is not sufficient, and when it exceeds 0.1%, the amount of oxygen in the molten steel is insufficient to form Ti-Mg complex oxide. This is because the amount decreases.

As described above, in the present invention, it is preferable to deoxidize the dissolved oxygen amount to 50-200 ppm before the addition of Ti. If the dissolved oxygen amount is less than 50 ppm, the amount of oxygen in molten steel is too small to form an appropriate amount of Ti-Mg oxide in order to show the effect of the present invention. There is a problem of oxidizing the alloying element to obtain an alloy component system composed of accurate chemical components.

As described above, the amount of dissolved oxygen is controlled using elements of Si, Mn, and Al, and then Mg and Ti are added as deoxidation elements. Since Mg has a high affinity with oxygen in molten steel, Mg reacts with oxygen present in molten steel to form Mg oxide. In addition, the Mg oxide has a high affinity with Ti, which has a direct effect on the formation of Ti-Mg composite oxide, which has a significant effect on increasing the number of Ti-based oxides distributed in the steel during solidification. When Ti and Mg are added at the same time, Mg has a greater affinity for oxygen than Ti and preferentially combines with oxygen in molten steel to form Mg oxides rather than Ti oxides. Therefore, since there is a problem of greatly reducing the number of Ti oxides acting as a polygonal or needle-like ferrite transformation nucleus, the ferrite nucleation is influenced by first distributing the Mg oxide in the deoxidation process by Mg and then deoxidizing by Ti. It is advantageous to increase the number of Ti oxides which affect the

In the present invention, Mg is added in the range of 0.001-0.005% as the tertiary deoxidation element, because the Mg content is less than 0.001%, the number of Mg oxides formed in molten steel is too small to form Ti-Mg complex oxides. This is because if the affinity between Mg and oxygen is greater than Ti oxide, the amount of oxygen that can be combined with Ti is insufficient at 0.005% or more, so that the number of TiO oxides as well as Ti-Mg composite oxides for showing the effect of the present invention is reduced. .

As described above, after the addition of the Mg deoxidation element, 0.005-0.2% Ti is added in the molten steel state. Rather than forming Ti oxide in the molten steel state, the solid solution in most molten steel and forming Ti oxide during solidification has an effect of increasing the number and finer Ti oxide. Therefore, when Ti is added after the addition of Mg deoxidation element, the ratio of remaining as solid solution Ti is rather high than that of oxides, thereby securing a fine amount of Ti oxide and TiN precipitates during solidification. At this time, the preferred amount of Ti is in the range of 0.005-0.2%, which considers the effective distribution of Ti oxide and TiN precipitates to show the effects of the present invention, as well as MnS, CuS, BN, and VN precipitation behaviors effective for improving the toughness of the weld heat affected zone. Range.

As described above, Mg is added followed by Ti, followed by refining, for example, degassing, wherein the holding time of the molten steel until continuous casting is preferably within 0.5-20 minutes. If the retention time of molten steel is less than 0.5 minutes, there is no time for coarse primary oxides to separate and float, and after solidification, the average size of Ti-based oxides is increased. It is very likely to be talked about and injured. Therefore, when Ti element is added within 0.5-20 minutes after Mg element is added, a fine Ti-Mg composite oxide can be obtained during solidification.

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. Below 1100 ° C., the number of TiN precipitates is small because the rate of diffusion of solute atoms is small. When it exceeds 1250 ° C., Ti-based precipitates are coarsened or decomposed, and thus the number of precipitates is reduced. I can't. 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 to the ferrite transformation end temperature ± 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 steel material 20㎛ or less.

[Welding Structure]

The welded steel material provided according to the present invention is a composite structure of ferrite (70% or more) + pearlite, and the size of the ferrite grains is 20 µm or less. In addition, TiN precipitates are 1.0 × 10 ⁷ or more per 1 mm ^{2 with} a size of 0.01-0.1 μm, and the interval is 0.5 μm or less, and Ti-Mg oxide is 1.0 × ¹⁰ per 1 mm ^{2 with} a size of 0.5-2.0 μm. ^It is finely distributed in the range of ² -1.0x10 ³ .

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

When the high heat input welding is applied and quenched as described above, the ferrite having the microstructure size of the heat-affected portion of about 20 μm or less has a phase fraction of about 70% or more. If the grain size of the ferrite is larger than 20 μm, the ferrite fraction of the side plate or allotriomorphs harmful to the weld heat affected zone toughness increases. When the ferrite of the present invention has the characteristics of polygonal ferrite and acicular ferrite, it is more advantageous for toughness, which is a Ti-Mg based composite oxide according to the present invention.

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. The slab may be subjected to hot charge rolling and direct rolling according to the user's use in the hot rolling process, and various known 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 in the slab by the continuous casting method by dissolving them in a converter, and the composition ratio between alloying elements by steel type to show the effect of the present invention is shown in Table 3. The solidification rate, slab heating temperature, heating time, rolling start temperature and end temperature, rolling reduction, and cooling rate of the rolled material having a thickness of 25 to 40 mm in the rolling process are shown in Tables 2 and 4. At this time, the rolling reduction ratio of all the steel grades was 60% 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 Mg 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 - 0.003 - 32 Inventive Steel 2 0.07 0.12 1.50 0.006 0.005 0.07 0.05 10 275 0.002 - 0.2 - - - 0.01 - 0.002 - 54 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 - 0.002 - 28 Inventive Steel 4 0.10 0.12 1.48 0.006 0.005 0.02 0.02 5 80 0.001 - - - - - 0.05 - 0.001 - 31 Inventive Steel 5 0.08 0.15 1.52 0.006 0.004 0.09 0.05 15 300 0.002 0.1 - 0.1 - - 0.05 - 0.002 - 52 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 - 0.001 - 32 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 - 0.001 - 30 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 - 0.002 - 44 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 0.002 - 43 Inventive Steel 10 0.07 0.16 1.45 0.008 0.006 0.045 0.025 6 100 0.05 - 0.3 - - 0.01 0.02 - 0.002 0.01 52 Inventive Steel 11 0.09 0.12 1.48 0.006 0.003 0.048 0.019 10 130 0.01 - 0.2 - - - - - 0.001 - 32 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 - 0.0022 - 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 - 0.0024 - 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 - 0.0019 - 25 Conventional Steel 7 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 - 0.3 - - - 0.036 - - - - - Conventional Steel 8 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 - 0.32 0.35 - - 0.013 - - - - - Conventional Steel 9 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 - - - 0.21 0.19 0.025 0.035 - - - - Conventional Steel 10 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 - - 1.09 0.51 0.36 0.021 0.021 - - - - Conventional Steel 11 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 - - - - - - 0.069 - - - - Conventional steels (1, 2, 3) are invention steels (5, 32, 55) of Japanese Patent Application Laid-Open No. 9-194990. Conventional steels (4, 5, 6) are Japanese Patent Application Laid-open Nos. Invented steels (14, 24, 28) and conventional steels (7, 8, 9, 10) are invented steels (48, 58, 60, 61) of Japanese Patent Application Laid-Open No. 8-60292. Is invention steel F of JP-A-11-140582

Steel grade used division Primary deoxidation sequence Amount of dissolved oxygen after Al deoxidation (ppm) Mg addition after Al deoxidation (%) Ti content after Mg addition (%) 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 68 0.003 0.014 15 30 1.1 0.35 Invention 2 Mn → Si 72 0.003 0.014 15 32 1.1 0.35 Invention 3 Mn → Si 56 0.003 0.014 15 34 1.1 0.35 Comparative Material 1 Mn → Si 35 0.003 0.014 15 9 1.1 0.35 Comparative Material 2 Mn → Si 264 0.003 0.014 15 134 1.1 0.35 Inventive Steel 2 Invention 4 Mn → Si 64 0.002 0.05 15 54 1.0 0.32 Invention Steel 3 Invention 5 Mn → Si 62 0.002 0.015 13 28 1.0 0.32 Inventive Steel 4 Invention 6 Mn → Si 62 0.001 0.02 14 31 1.0 0.35 Inventive Steel 5 Invention 7 Mn → Si 59 0.002 0.05 18 52 1.0 0.32 Inventive Steel 6 Invention Material 8 Mn → Si 58 0.001 0.02 15 32 1.0 0.35 Inventive Steel 7 Invention 9 Mn → Si 62 0.001 0.015 14 30 1.0 0.32 Inventive Steel 8 Invention 10 Mn → Si 73 0.002 0.018 16 44 1.1 0.34 Inventive Steel 9 Invention 11 Mn → Si 82 0.002 0.02 15 43 1.1 0.35 Inventive Steel 10 Invention Material12 Mn → Si 69 0.002 0.025 15 52 1.05 0.35 Inventive Steel 11 Invention Material 13 Mn → Si 62 0.001 0.019 15 32 1.1 0.32 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 Mg / O (Ti + Mg) / O Ti / N N / B Al / N V / N (Ti + 2Al + 4B + V) / N Invention 1 4.7 1.0 5.7 1.2 16.4 3.5 0.9 9.3 Invention 2 4.4 0.9 5.3 1.2 16.4 3.5 0.9 9.3 Invention 3 4.0 0.9 5.0 1.2 16.4 3.5 0.9 9.3 Comparative Material 1 15.6 3.3 18.9 1.2 16.4 3.5 0.9 9.3 Comparative Material 2 1.0 0.2 1.3 1.2 16.4 3.5 0.9 9.3 Invention 4 9.3 0.4 13.0 1.8 27.5 2.5 0.4 7.4 Invention 5 5.4 0.7 9.6 1.3 37.3 5.4 1.8 14.0 Invention 6 6.5 0.3 6.1 2.5 16.0 2.5 6.3 14.0 Invention 7 9.6 0.4 6.8 1.7 20.0 3.0 1.7 9.5 Invention Material 8 6.3 0.3 10.0 2.0 10.0 2.5 9.0 16.4 Invention 9 5.0 0.3 5.3 1.3 14.4 3.5 1.7 10.3 Invention 10 4.1 0.5 4.5 1.5 12.0 5.0 0.8 12.7 Invention 11 4.7 0.5 5.1 2.2 22.5 2.8 2.2 10.2 Invention Material12 4.8 0.4 5.2 2.5 16.7 4.5 2.0 13.7 Invention Material 13 5.9 0.3 6.3 1.5 13.0 3.6 - 9.0 Conventional Steel 1 0.4 - 0.4 4.1 13.8 0.6 - 5.7 Conventional Steel 2 3.8 - 3.8 2.5 96.0 0.8 - 4.0 Conventional Steel 3 0.7 - 0.7 0.8 105.8 0.4 - 1.5 Conventional Steel 4 5.2 0.9 6.08 4.1 2.5 0.8 8.8 24.5 Conventional Steel 5 4.8 0.9 5.7 6.5 0.25 10.5 18.5 4.7 Conventional Steel 6 3.6 0.8 4.36 3.2 0.39 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 880 65/80 7 550 Inventive Example 2 1200 130 1040 890 65/80 7 550 Inventive Example 3 1240 90 1040 880 65/80 7 550 Comparative Example 1 900 50 1040 880 65/80 7 550 Comparative Example 2 1350 250 1035 880 65/80 7 550 Invention 1 Inventive Example 4 1200 130 1020 890 65/80 6 550 Invention 3 Inventive Example 5 1200 130 1040 890 65/80 6 550 Comparative Material 1 Comparative Example 3 1210 120 1030 880 65/80 0.1 Room temperature Comparative Material 2 Comparative Example 4 1210 120 1030 890 65/80 19 Room temperature Invention 4 Inventive Example 6 1180 150 1020 890 60/80 7 550 Invention 5 Inventive Example 7 1190 140 1010 840 60/80 8 550 Invention 6 Inventive Example 8 1220 110 1010 810 60/75 7 550 Invention 7 Inventive Example 9 1190 130 1020 800 60/75 10 550 Invention Material 8 Inventive Example 10 1210 120 1010 890 60/75 10 550 Invention 9 Inventive Example 11 1220 110 1000 880 55/70 10 550 Invention 10 Inventive Example 12 1210 120 1010 890 55/70 9 550 Invention 11 Inventive Example 13 1180 130 1000 820 55/70 8 550 Invention Material12 Inventive Example 14 1210 110 1020 880 55/70 10 550 Invention Material 13 Inventive Example 15 1200 130 1020 880 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 2.4 X ^{10 8} 0.016 0.25 2.3 X 10 ² 1.8 25 394 553 38 11 74 358 Inventive Example 2 3.2 X ^{10 8} 0.017 0.24 2.1X10 ² 1.4 25 395 551 39 9 73 362 Inventive Example 3 2.5 X ^{10 8} 0.012 0.26 2.3 X 10 ² 1.2 25 396 550 39 10 73 357 Comparative Example 1 2.3 X 10 ⁶ 0.174 1.6 2.7 X 10 ¹ 2.2 25 393 554 26 16 50 206 Comparative Example 2 3.4 X 10 ⁶ 0.165 1.8 2.6 X 10 ¹ 2.3 25 792 860 17 17 21 45 Inventive Example 4 3.2 X ^{10 8} 0.025 0.32 3.2 X 10 ² 1.2 30 396 558 38 11 73 349 Inventive Example 5 2.6 X ^{10 8} 0.013 0.34 2.6 X 10 ² 1.2 30 396 562 38 10 73 354 Comparative Example 3 1.3 X 10 ⁶ 0.182 1.2 1.5X10 ¹ 3.4 30 384 564 30 18 53 220 Comparative Example 4 4.3X10 ⁶ 0.177 1.4 1.2X10 ¹ 3.7 30 392 582 29 17 64 208 Inventive Example 6 3.3 X ^{10 8} 0.026 0.35 1.8 X 10 ² 1.3 30 390 563 38 10 74 364 Inventive Example 7 4.6 X ^{10 8} 0.024 0.32 3.2 X 10 ² 1.2 35 390 564 39 10 75 360 Inventive Example 8 4.3X10 ⁸ 0.014 0.40 2.1X10 ² 1.6 35 392 542 36 11 76 365 Inventive Example 9 5.6 X ^{10 8} 0.028 0.29 2.4 X 10 ² 1.5 35 391 536 37 10 74 359 Inventive Example 10 5.2 X ^{10 8} 0.021 0.28 2.3 X 10 ² 1.5 35 394 566 36 10 73 375 Inventive Example 11 3.7 X ^{10 8} 0.029 0.25 2.7 X 10 ² 1.7 40 390 566 37 12 73 364 Inventive Example 12 3.2 X ^{10 8} 0.025 0.31 1.7 X 10 ² 1.5 40 396 542 38 11 75 356 Inventive Example 13 3.3 X ^{10 8} 0.042 0.34 1.8 X 10 ² 1.5 40 406 564 38 12 76 348 Inventive Example 14 3.6X10 ⁸ 0032 0.28 2.2 X 10 ² 1.6 40 387 550 37 10 74 349 Inventive Example 15 4.2 X ^{10 8} 0.018 0.26 1.9 X 10 ² 1.7 30 389 549 39 9 76 368 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 2.4X10 ⁸ / mm ² or more, whereas in the case of conventional steel 4.07 X10 ⁶ / mm ² Since the following ranges show that the invention material has a fairly uniform and fine precipitate size compared to the conventional material, the number is also significantly increased. In addition, the present invention has a number of oxides of Ti-Mg is also about 1.7x10 ² ~ 3.2x10 ³ ranges and the average size is about 1.2-1.8㎛. On the other hand, in the structure of the base steel of the present invention, in the case of the present invention, the ferrite grain size (FGS) is in the range of about 9-12 μm, which is very fine compared to that of the comparative material. It consists of fractions.

division Austenitic grain size of welding heat affected zone (㎛) Microstructure with welding heat effect of 100kJ / cm heat input Reproduction Weld Heat Affected Zone -40 ℃ Impact Toughness (J) (Maximum Heating Temperature: 1400 ℃) 1200 (℃) 1300 (℃) 1400 (℃) Ferrite Percentage (%) Ferrite Average Grain Size (㎛) Δt _800-500 = 60 seconds Δt _800-500 = 120 seconds Δt _800-500 = 180 seconds Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Inventive Example 1 23 34 56 74 15 372 -74 332 -67 293 -63 Inventive Example 2 22 35 55 77 13 384 -76 350 -69 302 -64 Inventive Example 3 23 35 56 75 13 366 -72 330 -67 295 -63 Comparative Example 1 54 86 182 38 24 124 -43 43 -34 28 -28 Comparative Example 2 65 92 198 36 26 102 -40 30 -32 17 -25 Inventive Example 4 25 38 63 76 14 353 -71 328 -68 284 -65 Inventive Example 5 26 41 57 78 15 365 -71 334 -67 295 -62 Comparative Example 3 56 80 178 40 26 108 -39 56 -32 24 -24 Comparative Example 4 63 88 184 39 28 64 -28 39 -30 10 -21 Inventive Example 6 25 32 53 75 14 383 -73 354 -69 303 -63 Inventive Example 7 24 35 55 77 14 365 -71 337 -67 292 -63 Inventive Example 8 27 37 53 74 13 362 -71 339 -67 296 -62 Inventive Example 9 24 36 52 78 15 368 -72 330 -67 284 -63 Inventive Example 10 22 34 53 75 14 383 -72 345 -66 293 -63 Inventive Example 11 26 35 64 75 14 356 -71 328 -68 282 -68 Inventive Example 12 27 39 64 74 15 353 -71 321 -67 276 -62 Inventive Example 13 23 38 68 74 14 354 -71 320 -67 254 -62 Inventive Example 14 25 35 64 70 15 342 -71 326 -67 248 -63 Inventive Example 15 23 36 53 76 16 349 -72 332 -68 293 -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 Ti-Mg composite 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 secondary austenite grains and promoting needle-like ferrite transformation within the grains, it is possible to provide a structural structural steel for welding that can simultaneously secure excellent heat input heat-affected zone toughness.

Claims (11)

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: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 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, 0.2≤Mg / O≤3, 3≤ (Ti + Mg) / O≤12 The welded structural steel having TiN precipitate and Mg-Ti composite oxide, which satisfies and is composed of the remaining Fe and other impurities, and has a microstructure of a composite structure of ferrite and perlite having a microstructure of 20 µm or less. 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. Steel for welded structures having a TiN precipitate and a composite oxide of Mg-Ti, characterized in that it satisfies. The steel material according to claim 1 or 2, wherein the steel comprises Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%. For the weld structure having a TiN precipitate and a composite oxide of Mg-Ti, characterized in that one or two or more selected, and one or two selected from the group of Ca: 0.0005-0.005%, REM: 0.005 to 0.05%. Steel. The steel material of claim 1, wherein the TiN precipitates of 0.01-0.1 μm are distributed at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less, and the Ti-Mg composite oxide of 0.5-2.0 μm is 1 × 10 ² -1. Welded structural steels with a composite oxide of TiN precipitates and Mg-Ti, characterized in that × 10 ³ pieces / mm ² distribution. 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: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 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, 0.2≤Mg / O≤3, 3≤ (Ti + Mg) / O≤12 After heating the slab composed of the remaining Fe and other impurities in the range of 1000-1250 ° C. for 60-180 minutes, hot rolling is performed at a rolling ratio of 40% or more in the austenite recrystallization zone, and then the ferrite transformation end temperature is ± 10 ° C. A method of manufacturing a welded structural steel having a TiN precipitate and a composite oxide of Mg-Ti, comprising cooling at a rate of 1 ° C / min or more. 6. The steel material according to claim 5, wherein V is contained in an amount of 0.01 to 0.2%, a ratio of V and N (V / N) is 0.3 to 9, and 7≤ (Ti + 2Al + 4B + V) / N≤17. Method for producing a welded structural steel having a composite oxide of TiN precipitate and Mg-Ti characterized in that it satisfies. The steel material according to claim 5 or 6, 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%. For the weld structure having a TiN precipitate and a composite oxide of Mg-Ti, characterized in that one or two or more selected, and one or two selected from the group of Ca: 0.0005-0.005%, REM: 0.005 to 0.05%. Method of manufacturing steels. According to claim 5, The slab is deoxidized element selected from the group of Al, REM, Zr, Ca, Mg having a higher deoxidizing power than Ti to deoxidize the dissolved oxygen amount to 50 ~ 200ppm before adding Mg, Ti, TiN precipitate and Mg-Ti composite, characterized in that the amount of Mg by adding Mg to 0.005-0.1%, and the amount of Ti by adding Ti to 0.005 to 0.2%, followed by degassing and continuous casting Method for producing a welded structural steel having an oxide. The method of claim 5 or 8, wherein the continuous casting of the slab is characterized in that the molten steel is cast at a rate of 0.9 ~ 1.2m / min and coldly cooled to a non-amount of 0.3 ~ 0.35ℓ / kg in the secondary cooling zone Method for producing welded steel having TiN precipitate and composite oxide of Mg-Ti. The heat input welding is applied to the steel material (base material) of claim 4, thereby producing a prior austenite of 80 µm or less in the weld heat affected zone, and then quenching the ferrite having a microstructure of 20 µm or less in the weld heat affected zone. Welded structure with excellent toughness of weld heat affected zone with over 70% phase fraction. The weld structure of claim 10, wherein the ferrite is polygonal ferrite and acicular ferrite.