KR20120014413A

KR20120014413A - High strength steel plate for welding structure with superior haz toughness for high heat input welding and method for manufacturing the same

Info

Publication number: KR20120014413A
Application number: KR1020100076457A
Authority: KR
Inventors: 정홍철
Original assignee: 주식회사 포스코
Priority date: 2010-08-09
Filing date: 2010-08-09
Publication date: 2012-02-17

Abstract

The present invention relates to steel materials for high strength welded structures used in welded structures such as ships, buildings, bridges, offshore structures, steel pipes, line pipes, etc. The object of the present invention is the welding heat effect during high heat welding using TiN and AlN stable at high temperatures. Fine TiN precipitates inhibit austenite grain growth during negative austenite grain growth, and the intragranular ferrite plate (IFF) structure with excellent AlN precipitates in the grains ensures the toughness of the high heat input welding heat affected zone. It is to provide an excellent high strength welded structural steel and its manufacturing method.
The present invention for achieving the above object, in the weight% C: 0.03-0.20%, Si: 0.01-0.5%, Mn: 0.6-2.0%, Ti: 0.005-0.2%, Al: 0.2-1.0%, B: 0.0005 to 0.01%, N: 0.008 to 0.03%, P: 0.03% or less (excluding 0), S: 0.03% or less (excluding 0), O: 0.03% or less (excluding 0), 1.2≤Ti / N Satisfies ≤2.5, 2≤N / B≤20, 20≤Al / N≤80, 40≤ (Ti + 2Al + 4B) / N≤200, is composed of the remaining Fe and other impurities, and the steel microstructure is ferrite The technical gist of the high-strength welded structural steel having excellent heat input weld heat-affected zone toughness.

Description

High Strength Welded Structural Steels with High Heat Resistance and High-Strength Weld Heat Affected Zones

The present invention relates to steel materials used in welding structures such as ships, buildings, bridges, offshore structures, steel pipes, line pipes, and more. It relates to a welded structural steel and a method of manufacturing the same.

In recent years, as buildings and structures become higher and larger, steel materials used in these buildings and structures are also being enlarged and their thickness is getting thicker. In order to weld such large thick steels, higher efficiency welding is required, and the heat input welding is generally used for large thick steels because it is advantageous in terms of construction cost reduction and welding construction efficiency. High heat input welding includes electro-gas welding and electro-slag welding, which can be used for 1 pass welding, and the current range of high heat input used is about 100 ~ 200kJ / cm, and the thicker plate thickness 50mm ~ 100mm is welded. In order to achieve this, the heat input range should be about 300 ~ 600kJ / cm.

In performing the above-mentioned high heat input welding, the base material side number mm is larger than the heat affected zone (bond part) of the weld-affected zone which is affected by the heat of the welding base material (steel plate as the weld material). The toughness in the position of?) Becomes a problem. In particular, the coarse grain HAZ near the fusion boundary is heated to a temperature close to the melting point by the heat input of the weld, so that the grain grows and the cooling rate decreases due to the increase in the heat input. Since the structure is easy to form and the microstructures that are vulnerable to toughness such as bainite and island martensite are formed during the cooling process, the toughness of the weld heat affected portion of the weld portion is likely to deteriorate.

Steels used in buildings and structures are required to have good toughness of welds as well as the strength of the base metal in terms of securing safety. Therefore, in order to secure the stability of welded structures, the growth of austenite grains in the weld heat affected zones is suppressed to be fine. It needs to be maintained. As a means to solve this problem, a technique for delaying grain growth of the weld heat affected zone during welding by appropriately dispersing oxides or Ti-based carbonitrides stable at high temperatures in steel materials has been described. 298708, (Pyeong) 10-298706, (Pyeong) 9-194990, (Pyeong) 9-324238, (Pyeong) 8-60292, (S) 60-245768, (Pyeong) 5-186848, (S) 58- 31065, (S) 61-79745, Japanese Journal of Welding Society 52, 2, page 49 and Japanese Patent Application Laid-Open No. 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 kJ / cm of heat input (maximum heating temperature of 1400 ° C.) is applied. Is about 300J). In the patent 0.05㎛ than TiN precipitates to substantially manage to 4-12 the Ti / N is 5.8 × 10 ³ ³ /㎟~8.1×10 gae gae / ㎟, In addition, 0.03 ~ 0.2㎛ of TiN precipitates is 3.9 × 0 The toughness of the welded portion is secured by making the ferrite fine by depositing ³ pieces / mm 2 to 6.2 × 10 ⁴ / mm 2.

However, according to the above patent, 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 with the highest impact toughness of 0 ° C.) and the base material and heat affected zone. As the negative toughness difference is about 100J, there is a limit in securing the reliability of steel structure by the heat input welding of thickened steel. In addition to this, as a method for securing the desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or higher and quenched and then reheated for hot rolling. There is.

In addition, Japanese Laid-Open Patent Publication No. 9-194990 manages the ratio of Al and O to 0.3 ≦ Al / O ≦ 1.5 in low nitrogen steels (N ≦ 0.005%) to control composite oxides of Al, Mn, and Si. When the technology used or the high heat input welding of about 100 kJ / cm is applied, the welding heat affected zone transition temperature is -50 ° C, and the toughness is not good, and Japanese Patent Application Laid-Open No. 10-298708 is a MgO-TiN composite. When a technique using a precipitate or a large input heat welding of about 100 kJ / cm is applied, the impact toughness of the welding heat affected zone 0 ° C. is 130 J, which is inferior in toughness.

Until now, many techniques have been known to improve the toughness of the weld heat affected zone by using TiN precipitates and Al or Mg oxides during high heat input welding. There is no technology that has been significantly improved. Therefore, if the welding heat-affecting zone can have the same level of toughness as the base metal, the heat input welding is possible for large thick steels such as buildings and structures, and the welding heat-affecting zone, which can secure stability and reliability, is equivalent to that of the base metal. It is time to develop steel that can have toughness.

One aspect of the present invention is to provide a high-strength welded structural steel material and a method of manufacturing the same, by uniformly dispersing TiN and AlN precipitates in the steel during the production of high-strength welded steel structure, the toughness of the weld heat affected zone during welding.

In the present invention, C: 0.03 to 0.2%, Si: 0.01 to 0.5%, Mn: 0.6 to 2.0%, Ti: 0.005 to 0.2%, Al: 0.2 to 1.0%, B: 0.0005 to 0.01%, and N: Heat input welding heat composed of 0.008 ~ 0.030%, P: 0.03% or less (excluding 0), S: 0.03% or less (excluding 0), O: 0.03% or less (excluding 0), remaining Fe and other impurities Provide high strength welded structural steel with excellent impact toughness.

By weight%, C: 0.03-0.23%, Si: 0.01-0.5%, Mn: 0.6-2.0%, Ti: 0.005-0.2%, Al: 0.2-1.0%, B: 0.0005-0.01%, N: 0.008-0.030 %, P: 0.03% or less (excluding 0), S: 0.03% or less (excluding 0), O: 0.03% or less (excluding 0), slabs composed of remaining Fe and other impurities at 1000 ~ 1250 ℃ A heating step of heating for 120 to 240 minutes; A hot rolling step of hot rolling the heated slab at a rolling ratio of at least 40% in an austenite recrystallization zone; And a cooling step of cooling the hot rolled steel sheet at a temperature of 10 ° C./sec or less from a rolling end temperature to a room temperature.

The present invention is to develop a high-strength welded structural steel material that can secure the properties of high heat input welding heat affected zone while having a high-strength base material properties at the same time to enable more efficient welding, reducing the construction cost and improve the efficiency of welding construction It works. In addition, there is an effect that enables large heat input welding in a state in which stability and reliability for large thick steels are secured.

In the present invention, in order to improve the weld toughness, the size of the weld heat-affected zone austenite grains is 80 micrometers or less during the high heat input welding using precipitates in steel, and the addition of an appropriate amount of aluminum promotes ferrite transformation and also crystallization. AlN precipitate is used in the mouth to promote needle ferrite transformation.

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.

The inventors of the present invention have found that the toughness of the weld heat affected zone is significantly changed based on the grain size of about 80 μm of the austenite, and the grain size of the old austenite is determined by the size and number of precipitates. Focusing on the distribution, the size of the old austenite grains was less than 80 µm using TiN precipitates.

TiN precipitates precipitated in steel have the effect of suppressing grain growth of austenite during high heat input welding. When high heat input welding is applied to structural steel, the weld heat affected zone is heated to a high temperature of about 1400 ℃ or more and TiN precipitates are decomposed. Or some of the precipitates are coarsened so that the grain growth inhibiting effect of the austenite as described above is eliminated, and thus the size of the austenite grains is coarse to 80 µm or more, resulting in a decrease in toughness of the weld heat affected zone.

When the high heat input welding is applied, the TiN precipitates in the steel are decomposed and the decomposed Ti atoms diffuse to cause the decomposition and coarsening of the TiN precipitates. In the present invention, the high temperature stability of the TiN precipitates is found in a high nitrogen environment. By using this, TiN precipitates were reduced in decomposition and coarsening. That is, in a high nitrogen environment (low Ti / N ratio), the solid solution Ti concentration and the diffusion rate of the solid solution Ti atoms are reduced to improve the high temperature stability of the TiN precipitates. In particular, when the ratio of Ti and N (Ti / N) is in the range of 1.2 to 2.5, the amount of Ti dissolved is extremely reduced, and the high temperature stability of the TiN precipitate is greatly improved, so that the fine TiN precipitate having a size of 0.01 to 0.1 μm is obtained. The result was distributed at 1.0 × 10 ⁷ holes / mm 2 or more at intervals of 0.5 μm or less.

In addition, in the present invention, by adding an appropriate amount of aluminum to promote the diffusion of carbon to prevent phase martensite transformation deterioration of impact toughness and to promote ferrite transformation.

According to the results of the present inventors, in general, the aluminum element has a very low solubility of 0.7% in the austenite phase, but the solubility of the ferrite phase is increased by 50 times to 30%. Therefore, if an appropriate amount of aluminum is added to the steel, carbon is pushed out of the ferrite phase while being heated to about 1400 ° C or more during welding and then transformed from the austenite phase to the ferrite phase during the cooling process while the aluminum element is dissolved in the ferrite phase. Is concentrated on the retained austenite and prevents transformation into MA Constituent and promotes the formation of carbides such as Fe ₃ C to promote pearlite transformation. That is, a large amount of aluminum is added to promote the diffusion of carbon to prevent phase martensite transformation, which degrades impact toughness, thereby improving impact toughness of the weld heat affected zone.

In addition, AlN precipitates are used to promote the transformation of intragranular ferrite plates, which promotes ferrite transformation, which is advantageous for toughness, and further improves the ferrite lath width and affects the fracture behavior, thereby greatly improving the weld heat affected zone toughness. .

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

First, the steel component will be described. The content of each element below represents weight percent.

It is preferable to make content of carbon (C) into 0.03 to 0.20%.

Carbon (C) is required 0.03% or more to secure the strength of the steel and weld heat affected zone. However, if the content exceeds 0.20%, the strength increases, but cracking occurs during welding, which is not preferable because it damages the weldability.

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

If the silicon content is less than 0.01%, the deoxidation effect of the molten steel is insufficient during steelmaking and the corrosion resistance of the steel is deteriorated. Therefore, it is preferable to add more than 0.01%, but if the content exceeds 0.5%, the effect is saturated and after rolling It is not preferable because it promotes transformation of the island martensite with increasing hardenability during cooling, thereby lowering the low temperature impact toughness and affecting the weld cracking susceptibility.

The content of manganese (Mn) is preferably limited to 0.6 ~ 2.0%.

Mn is an indispensable element to secure the strength and toughness of the steel, but it is necessary to add more than 0.6%, but when it contains more than 2.0%, the weld hardness of the weld heat affected zone during welding increases the weld cracking and welding It is not preferable because it lowers the toughness of the heat affected zone.

The content of aluminum (Al) is preferably limited to 0.2 ~ 1.0%.

Al is one of the most important elements, which promotes carbon diffusion when austenite-to-ferrite is transformed to prevent the formation of MA constituent phases, and also forms AlN in combination with N in the grains to promote needle-ferrite transformation. It plays a role. Therefore, in order to play the role, an aluminum amount of 0.2% or more is required, and when added in excess of 1.0%, coarse ferrite and martensite are mixed to greatly reduce the impact toughness of the weld heat affected zone.

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

Ti combines with N to form stable fine TiN precipitates at high temperature, thereby inhibiting austenite grain growth of the weld heat affected zone. In order to obtain a fine TiN precipitate dispersion effect, it is preferable to add more than 0.005% of Ti, and when it is added in excess of 0.2%, coarse TiN precipitate is formed in molten steel and mixed into the slab and steel material for welding. It is not preferable because minor austenite grain growth is not inhibited.

The content of boron (boron, B) is preferably limited to 0.0005 ~ 0.01%.

B is an element that improves the hardenability and segregates at the weld heat-affected austenite grain boundary to combine with N to form a BN precipitate and promotes fine polygonal ferrite transformation at the grain boundary. Therefore, 0.0005% or more is required, and if it exceeds 0.01%, it is not preferable because the weld hardenability is greatly increased to promote martensite transformation, which causes low temperature cracking of the weld and lowers the toughness of the weld heat affected zone. Therefore, the B content is limited to 0.0005 to 0.01%.

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

N is an indispensable element for forming TiN precipitates and the like. In the case of high heat input welding, N inhibits the growth of the old austenite grains at the weld heat affected zone and increases the amount of precipitates such as TiN. 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 the toughness decreases due to the increase in the amount of solid solution nitrogen distributed in the weld heat affected zone. Therefore, N content is limited to 0.008 to 0.03%.

The content of phosphorus (P) and sulfur (S) is preferably limited to 0.03% or less (excluding 0), respectively.

P is preferably as low as possible because it is an impurity element that promotes central segregation during rolling and hot cracking during welding. To improve steel toughness, weld heat affected zone toughness and reduce center segregation, it is recommended to manage it to 0.03% or less (excluding 0).

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 steel toughness, weld heat affected zone toughness and center segregation, it is recommended that the S content be 0.03% or less (excluding 0).

The content of oxygen (O) is preferably limited to 0.03% or less (excluding 0).

O is an element that reacts with Ti to form Ti oxide, which promotes the transformation of acicular ferrite in the weld metal. When the O content exceeds 0.03%, coarse oxides and other oxides such as FeO are formed, which is not preferable because it affects the properties of the weld heat affected zone.

In addition to the above composition range, the present invention preferably satisfies the following formula.

It is preferable to make ratio of Ti / N into 1.2-2.5.

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 TiN precipitate is more stable when the nitrogen content is higher than when 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. On the other hand, when the Ti / N ratio is less than 1.2, the amount of solid solution nitrogen of the steel is increased, which is harmful to the toughness of the weld heat affected zone.

It is preferable to make ratio of N / B into 2-20.

In the present invention, when the N / B ratio is less than 2, the number of BN precipitates that suppress the growth of the austenite grains during the cooling process after welding is insufficient, and when the N / B ratio exceeds 20, the effect is saturated and the amount of solid solution nitrogen is increased. This is because the toughness of the weld heat affected zone is lowered.

It is preferable to make Al / N ratio into 20-80.

In the present invention, when the Al / N ratio is less than 20, the distribution of AlN precipitates that can promote acicular ferrite transformation in the grains is insufficient, and when the Al / N ratio exceeds 80, the effect is saturated.

It is preferable to make ratio of (Ti + 2Al + 4B) / N into 40-200.

In the present invention, when the ratio of (Ti + 2Al + 4B) / N is less than 40, the size and distribution number of TiN, AlN, (Ti-Al) N precipitates for inhibiting the growth of the austenite grain growth of the weld heat affected zone are insufficient. This is because the effect is saturated when (Ti + 2Al + 4B) / N exceeds 200.

In the present invention, in order to further improve the mechanical properties to the steel composition as described above, one or more selected from the group of Ni, Cu, Mo, Cr, V, W, Zr, Nb may be 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 steel by solid solution strengthening. In order to achieve this effect, the Ni content is preferably contained 0.1% or more, but when it exceeds 3.0%, the hardenability is greatly increased to reduce the toughness of the weld heat affected zone and the occurrence of high temperature crack in the weld heat affected zone and the weld metal. It is not preferable because there is possibility.

The content of copper (Cu) is preferably limited to 0.1 ~ 2.0%.

Cu is employed at the base and is an effective element to secure the steel strength and toughness due to the effect of solid solution strengthening. To this end, the Cu content should be contained 0.1% or more, but when it exceeds 2.0% is not preferable because it increases the hardenability in the weld heat affected zone to lower the toughness and promotes high temperature crack in the weld heat affected zone and the weld metal.

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.

It is preferable to make chromium (Cr) into 0.05 to 1.0%.

Cr increases the hardenability and also improves the strength. If the content thereof is less than 0.05%, the strength cannot be obtained, and if it exceeds 1.0%, the base metal and the weld heat affected zone toughness deteriorate.

Molybdenum (Mo) is preferably made 0.05 to 1.0%.

Mo is also an element that increases the hardenability and at the same time improves the strength. The content is 0.05% or more for securing the strength, but the upper limit is 1.0% as in the case of Cr to suppress the hardening of the weld heat affected zone and the occurrence of low temperature welding crack. do.

The content of V is preferably limited to 0.005 ~ 0.1%.

V is preferably added in an amount of 0.005% or more for precipitation hardening, and further, because of the effect of depositing VN or VC to suppress the growth of the former austenite grains in the weld heat affected zone. However, if the V content exceeds 0.1%, it is not preferable because it increases the low-temperature transformation structure in the weld heat affected zone, which adversely affects the mechanical properties and adversely affects the weld cracking susceptibility.

The content of W is preferably limited to 0.005 ~ 0.5% or less.

W is an effective element for improving high temperature strength and strengthening precipitation. However, less than 0.005% is not preferable because the strength increase effect is weak, and if it exceeds 0.5%, it is not preferable because it adversely affects the toughness of the weld heat affected zone.

The content of Nb is preferably limited to 0.005 to 0.2%.

More than 0.005% is required to anticipate the effect of strengthening precipitation and improving the toughness of steel. However, exceeding 0.2% is undesirable because it promotes the formation of phase martensite in the weld heat affected zone during welding, which adversely affects the toughness of the weld heat affected zone.

The content of Zr is preferably limited to 0.05 to 0.5%.

Since Zr is effective in increasing the strength, it is preferable to add 0.05% or more. If it exceeds 0.5%, Zr is not preferable because it adversely affects the toughness of the weld heat affected zone.

In addition, in the present invention, one or two kinds of Ca and REM may be further added to suppress grain growth of the austenite.

Ca and REM form oxides with excellent high temperature stability to suppress the growth of austenite grains when heated in steel and promote ferrite transformation during cooling to improve the toughness of the weld heat affected zone. In addition, Ca has an effect of controlling the formation of coarse MnS during steelmaking. To this end, Ca 0.0005% or more, REM is preferably added 0.005% or more, if Ca exceeds 0.005% or REM exceeds 0.05% to create large inclusions and clusters (cluster) 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.

In the present invention, after the hot rolling, the microstructure of the steel is ferrite and perlite structure, and the ferrite is composed of 70 to 75%, and the pearlite is 20 to 25% remaining bainite and martensite. The reason is that it is possible to have bainite structure or martensite structure to increase the strength of the steel, but in this case, the strength is high but the toughness of the steel is difficult to secure, so that both strength and impact toughness can be secured.

In addition, in order to control the size of the old austenite grains below 80 μm, the precipitates distributed in the steel have a size of 0.01 to 0.1 μm, the number of 1.0 × 10 ⁷ to 1.0 × 10 ¹² per ¹ mm 2, and the interval between the precipitates. It is desirable to have a distribution so as to be 0.5 탆 or less. If the particle size of TiN precipitate and AlN precipitate is less than 0.01㎛, it is easily reusable to steel materials during high heat input welding, so that the effect of inhibiting the growth of old austenite grains is inadequate. The size is limited to 0.01 ~ 0.1㎛ because pinning (grain growth suppression) effect is less and the behavior like coarse nonmetallic inclusions has a detrimental effect on mechanical properties. On the contrary, if the number of the precipitates not less than 1.0 × 10 ² per 1mm less than ^7, difficult to control the old austenite grain size is substituted for hot welding heat affected zone in the threshold 80㎛ 1.0 × 10 ^12, the precipitate is too large precipitates in the base material Since the presence of this has an adverse effect on the base metal impact toughness, the number of precipitates is limited as described above.

In addition, in order to suppress Oswald lifeening phenomenon in which some precipitates become coarse, the distribution of precipitates should be uniform with the interval of TiN and AlN precipitates of 0.5 μm or less.

Hereinafter, a method for manufacturing a welded structural steel of the present invention will be described in detail.

In general, the steel refining process consists of an out-of-furnace refining process after the first refining of the converter and the second refining of the molten steel of the converter by ladle. ). Usually deoxidation takes place between primary and secondary refining.

Next, it is preferable to add Ti so that the content of Ti is 0.005 to 0.2%, which is an advantageous range for proper TiN precipitate distribution for showing the effect of the present invention.

In the present invention, the molten steel refined as described above is continuously cast into slabs. Continuous casting is preferred to cast at low speed in consideration of the high possibility of occurrence of cast surface cracks in high nitrogen steel, and to impart weak cooling conditions in the secondary cooling zone. 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 more preferably about 0.9 to 1.3 m / min or less than 1.3 m / min, which is lower than the conventional casting speed of about 1.5 m / min. The reason is that the casting speed is less than 0.9m / min, it is advantageous for the surface crack, but the productivity is lowered. If it is faster than 1.3m / min, the surface crack is more likely to occur.

In addition, in the secondary cooling zone, the specific amount is preferably weak cooling, that is, 0.3 to 0.35 l / kg. If the specific water content is less than 0.3 L / kg, it is difficult to control the appropriate size and number of TiN and AlN precipitates to show the effect of the present invention by coarsening TiN precipitates. In addition, when the specific amount exceeds 0.35ℓ / kg, the precipitation frequency of TiN and AlN precipitates is small, it is difficult to control the number and size of TiN precipitates for showing the effect of the present invention.

In the present invention, the slab is heated for 120 to 240 minutes at 1000 ~ 1250 ℃. At less than 1000 ℃, the number of TiN and AlN precipitates is small because the rate of diffusion of solute atoms is small. If it exceeds 1250 ℃, the TiN and AlN precipitates are coarsened or partially decomposed, and the number of precipitates is reduced. Not desirable On the other hand, in less than 120 minutes of heating time, the diffusion effect of solute atoms is insufficient, and thus, the slab segregation reduction effect is small, and also it is not preferable because there is not enough time for diffusion of solute atoms to form precipitates. In addition, when the heating time exceeds 240 minutes coarsening of austenite grain size occurs and 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 in the austenite uncrystallized temperature range. The austenite recrystallization zone temperature is influenced by the emphasis and the previous reduction amount, and the austenite recrystallization zone temperature is about 800 to 1000 DEG C in consideration of the conventional reduction amount in the emphasis of the present invention.

After rolling as described above, the air is cooled to room temperature at the end of the final rolling or accelerated cooling at a rate of 10 ° C./sec or less to room temperature to secure fine ferrite. If the cooling rate is cooled to a cooling rate of 10 ℃ / sec or more is not preferable because the transformation to martensite structure instead of the fine ferrite adversely affects the physical properties of the steel.

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, so that 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 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 and end temperature, rolling reduction, and rolling rate of 25 ~ 50mm thickness in rolling process. At this time, the cumulative reduction ratio during rolling of all steel grades was 70% or more.

The test pieces for evaluating the mechanical properties of the steel from the hot rolled plates as described above were taken from the center of the plate thickness of the rolled material, the tensile test was measured in the rolling direction, and the Charpy impact test in the direction perpendicular to the rolling direction. It was.

Tensile test pieces were used in the KS standard (KS B 0801) No. 4 test pieces, the tensile test was tested at a cross head speed (5mm / mim). The impact test piece was manufactured in accordance with KS (KS B 0809) No. 3 test piece, in which the notch direction was processed on the side of the rolling direction (L-T) in the case of steel, 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 weld 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 a simulated welding simulator. After holding, it was quenched with He gas. The quenched test piece was ground and corroded to determine the austenite grain size under the highest heating temperature conditions according to the KS standard (KS D 0205), and the results are shown in Table 4.

After cooling, the size, number, and spacing of precipitates and oxides, which have a significant effect on the analysis of the microstructure and the toughness of the weld affected zone, were measured by the point counting method using an image analyzer and electron microscope. Is shown in Table 4. At this time, the test surface was evaluated based on 100 mm ² .

Impact toughness evaluation of the weld 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 given, the surface of the test piece was polished, processed into an impact test piece, and evaluated through a Charpy impact test at -40 ° C. The results are shown in Table 5.

Ti / N N / B Al / N (Ti + 2Al + 4B) / N Invention 1 1.2 3.1 35 73.8 Invention 2 1.8 9.2 33 67.7 Invention 3 1.3 3.4 54 161.9 Invention 4 2.5 3.2 31 66.3 Invention 5 1.7 12.0 30 62.0 Invention 6 2.0 3.3 25 53.2 Invention Material7 1.3 4.1 35 71.8 Invention Material 8 1.5 4.8 50 102.3 Invention Material 9 2.2 2.6 39 81.5 Invention 10 2.5 3.9 45 93.5 Invention Material 11 1.5 8.7 37 75.4 Conventional Materials 1 4.1 13.8 0.6 5.7 Conventional material 2 2.5 96.0 0.8 4.0 Conventional Materials 3 0.8 105.8 0.3 5.5 Conventional Materials 4 4.1 4.0 0.8 15.5 Conventional material 5 6.5 4.0 1.1 35.6 Conventional Materials 6 3.2 2.5 0.4 21.9 Conventional Materials 7 1.0 9.9 2.5 11.0 Conventional Materials 8 1.2 14.3 0.4 4.5 Conventional Materials 9 0.8 9.1 2.1 9.6 Conventional Materials 10 0.6 9.5 3.2 10.4 Conventional Materials 11 5.5 12.7 3.4 20.3

As shown in Table 4, the precipitates (TiN and AlN precipitates) of the hot rolled material produced by the present invention have a number of 2 × 10 ⁸ pieces / mm ² or more, a size of 0.029 μm or less and a space of 0.45 μm, The conventional materials have a number of 4.3 × 10 ⁸ / mm ² or more, a size of 0.155 μm or less and a space of 1.5 μm or less, indicating that the number of the inventions is considerably increased while having a fairly uniform and fine precipitate size compared to the conventional material. have.

Table 5 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 weld heat affected zone, has a value of 62 μm or less in the case of the present invention, and a very coarse range of about 185 μm or more in the conventional case. . 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 microstructure of the steel of the present invention in the weld heat affected zone corresponding to the welding heat input amount of 100kJ / cm is composed of needle-like ferrite and polygonal ferrite in the mouth, the phase fraction of the two phases is composed of more than 90%, and excellent heat input compared to conventional materials The impact toughness of the weld heat affected zone is shown.

Claims

By weight%, C: 0.03-0.2%, Si: 0.01-0.5%, Mn: 0.6-2.0%, Ti: 0.005-0.2%, Al: 0.2-1.0%, B: 0.0005-0.01%, N: 0.008-0.030 High heat resistance welded heat affected zone toughness including%, P: 0.03% or less (excluding 0), S: 0.03% or less (excluding 0), O: 0.03% or less (excluding 0), remaining Fe and other impurities This excellent high strength welded structural steel.

The method according to claim 1,
Ti, N, B, and Al satisfy a relationship of 1.2≤Ti / N≤2.5, 2≤N / B≤20, 20≤Al / N≤80, and 40≤ (Ti + 2Al + 4B) / N≤200 High strength welded structural steel with excellent toughness.

The method according to claim 1,
The steel is a high-strength welded structural steel with excellent toughness of heat input heat-affected zone having a microstructure composed of ferrite 70 ~ 75%, pearlite 20 ~ 25% and the balance bainite and martensite.

The method according to claim 1,
The steel is V: 0.005 to 0.1%, Nb: 0.005 to 0.2%, Ni: 0.1 to 3.0%, Cu: 0.1 to 2.0%, Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0%, W: 0.05 to 0.5 %, Zr: High strength welded structural steel with excellent toughness of high heat input welded heat affected zone containing one or more selected from the group consisting of 0.05% to 0.5%.

The method according to claim 1,
The steel material is a high strength welded structural steel with excellent toughness of the high heat input welding heat affected zone comprising one or more selected from the group consisting of Ca: 0.0005 ~ 0.005%, REM: 0.005 ~ 0.05%.

The method according to claim 1,
High-strength welded structural steel with excellent toughness of high heat input heat-affected zones in which TiN precipitates and AlN precipitates of 0.01 to 0.1 ㎛ size and 1.0 to ¹⁰ pieces / mm2 or more are distributed at intervals of 0.5 µm or less.

The method according to claim 1,
When the steel is subjected to high heat input welding, austenite (prior austenite) of 80 μm or less is formed in the weld heat affected zone, and when quenched, more than 90% of the microstructures of the weld heat affected zone are formed of in-mouth acicular ferrite and ferrite structure. High-strength welded structural steel with excellent toughness.

By weight%, C: 0.03-0.23%, Si: 0.01-0.5%, Mn: 0.6-2.0%, Ti: 0.005-0.2%, Al: 0.2-1.0%, B: 0.0005-0.01%, N: 0.008-0.030 %, P: 0.03% or less (excluding 0), S: 0.03% or less (excluding 0), O: 0.03% or less (excluding 0), slabs composed of remaining Fe and other impurities at 1000 ~ 1250 ℃ A heating step of heating for 120 to 240 minutes;
A hot rolling step of hot rolling the heated slab at a rolling ratio of at least 40% in an austenite recrystallization zone; And
The method of manufacturing a high strength welded structural steel having excellent toughness in heat input welding heat affected zone including a cooling step of cooling the hot rolled steel sheet at a rate of 10 ° C./sec or less from a rolling end temperature to room temperature.

The method according to claim 8,
Ti, N, B, and Al satisfy a relationship of 1.2≤Ti / N≤2.5, 2≤N / B≤20, 20≤Al / N≤80, and 40≤ (Ti + 2Al + 4B) / N≤200 High heat welded welded steel affected zone toughness high strength welded structural steel manufacturing method.

The method according to claim 8,
The steel is V: 0.005 to 0.1%, Nb: 0.005 to 0.2%, Ni: 0.1 to 3.0%, Cu: 0.1 to 2.0%, Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0%, W: 0.05 to 0.5 %, Zr: 0.05 to 0.5% high heat resistance weld heat affected zone comprising one or two or more selected from the group consisting of high strength welded structural steel manufacturing method.

The method according to claim 8 or 10,
Wherein the steel is Ca: 0.0005 ~ 0.005%, REM: 0.005 ~ 0.05% of the high heat resistance welded heat affected zone toughness, characterized in that it contains at least one selected from the group consisting of a group consisting of high strength welded structural steel.

The method according to any one of claims 8 to 10,
The slab is a high strength welded structure with excellent toughness of high heat input heat-affecting zone, characterized in that the molten steel is continuously cooled at a rate of 0.9 ~ 1.3m / min at a cold amount of 0.3 ~ 0.35ℓ / kg in the secondary cooling zone Method of manufacturing steels.

The method of claim 11,
The slab is a high strength welded structure with excellent toughness of high heat input heat-affecting zone, characterized in that the molten steel is continuously cooled at a rate of 0.9 ~ 1.3m / min at a cold amount of 0.3 ~ 0.35ℓ / kg in the secondary cooling zone Method of manufacturing steels.