KR20020014133A - Steel plate having superior toughness in weld heat-affected zone and method for manufacturing the same, welding fabric using the same - Google Patents

Abstract

PURPOSE: A steel product for a welding structure is provided which has superior toughness in a heat-affected zone as a difference between a base metal and the heat-affected zone is being minimized even in the weld heat input ranging from middle heat input to very large heat input using TiN and TiO oxides, and a method for manufacturing the same and a welding structure using the same are provided. CONSTITUTION: The method for manufacturing a steel product for a welding structure having superior toughness in a heat-affected zone comprises the processes of heating a steel slab in the temperature range of 1100 to 1250 deg.C for 60 to 180 minutes, a composition of the steel slab comprises 0.03 to 0.17 wt.% of C, 0.01 to 0.5 wt.% of Si, 0.4 to 2.0 wt.% of Mn, 0.005 to 0.2 wt.% of Ti, 0.0005 to 0.1 wt.% of Al, 0.008 to 0.030 wt.% of N, 0.0003 to 0.01 wt.% of B, 0.001 to 0.2 wt.% of W, 0.03 wt.% or less of P, 0.03 wt.% or less of S, 0.002 to 0.03 wt.% of O and a balance of Fe and other impurities, wherein 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; hot rolling the heated steel slab at a 40% or more of rolling ratio in the austenite recrystallization temperature zone; and cooling the hot rolled steel slab to a finishing temperature of ferrite transformation±10 deg.C in a cooling rate of 1 deg.C/min or more.

Description

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Recently, the steel used in accordance with the trend of high-rise building, structure has been replaced by thick steel.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

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

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

Therefore, in order to secure the stability of the welded structure, it is necessary to suppress the growth of the austenite grains in the weld heat affected zone and to keep it fine. As a means to solve this problem, there is disclosed a technique for delaying grain growth of the weld heat affected zone during welding by appropriately distributing an oxide or Ti-based carbonitride, which is stable at a high temperature, to steel materials. For example, Japanese Patent Application Laid-Open No. 11-140582, No. 10-298708, No. 10-298706, No. 9-194990, No. 9-324238, No. 8-60292, (S) 60-245768, (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 discloses a technique using MgO-TiN composite precipitates, but the impact toughness of the welding heat-affected zone at 0 ° C. is 130J when tough heat welding of about 100 kJ / cm is applied. Not good.

Until now, many techniques have been known to improve the toughness of weld heat affected zones during high heat input welding using TiN precipitates and Al-based or MgO oxides. No significant improvements have been made yet. In particular, there are few technologies in which the toughness of the weld heat affected zone shows an equivalent level to that of the base metal. Therefore, if the above problems can be solved, super heat input welding of the thickened steel is possible, so that it is possible to achieve high efficiency of the welding operation as well as to secure the high-rise of the steel structure and the reliability of the steel structure.

The present invention is a welded structural steel material having excellent toughness of the heat-affected portion and its manufacturing method, while the difference in the toughness of the base material and the heat-affected portion is minimized even in the weld heat input range from medium to high heat input using TiN and TiO oxide, and It is an object of the present invention to provide a welded structure using heat affected zone characteristics of structural steels.

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%, 1.2≤Ti / N≤2.5, Satisfying 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10, and are composed of the remaining Fe and other impurities, The microstructure is composed of a composite 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%, 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, and the steel slab is 1100 with the remaining Fe and other impurities After heating for 60-180 minutes in the range of -1250 ° C, hot rolling is carried out in the austenite recrystallization zone with a rolling ratio of 40% or more, and then cooling to a ferrite transformation end temperature ± 10 ° C at a rate of 1 ° C / min or more.

In addition, in the welded structure of the present invention, welding is applied to the welded structural steel (base metal) of the present invention to generate a prior austenite of 80 μm or less in the weld heat affected zone, and then quenched to weld heat affected. Negative microstructure is composed of ferrite with 20 micrometers or less in phase fraction 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, and the austenite formed in the manufacturing process of the steel (hot rolling process). Use for convenience to distinguish from.

As a result of investigating the old austenite grain size on the toughness of the weld heat affected zone, the inventors found that the toughness of the weld heat affected zone changes based on the grain size (about 80 μm) of the critical old austenite.

Based on these studies, in the present invention,

[1] with the use of TiN precipitates and TiO oxides,

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

[3] Effectively use TiO oxide, BN and AlN precipitates to increase the ferrite fraction in the weld heat affected zone, and enhance the toughness improvement effect by promoting the transformation of polygonal or needle-like ferrite in Guostenite. These [1] [2] [3] are demonstrated in more detail.

[1] TiN precipitate and TiO oxide management

When heat input welding is applied to structural steels, the weld heat affected zone near the melting line is heated to a high temperature of about 1400 ° C or more, and TiN precipitates precipitated in the base metal are partially dissolved by welding heat or Ostwald ripening. Small precipitates are decomposed and diffuse into larger precipitates, resulting in larger precipitates). Only some precipitates are coarse, and the number of TiN precipitates is significantly reduced, thereby suppressing the inhibitory effect of the growth of austenite grains. .

The inventors have found that this phenomenon is caused by the diffusion of TiN precipitates dispersed in the base metal by the dissolution of the dissolved Ti atoms by the heat of welding, and the characteristics of TiN precipitates according to the ratio of Ti / N are as follows. / N ratio is low), it is found that the dissolved Ti concentration and diffusion rate of the dissolved Ti atoms and the high temperature stability of the TiN precipitates are improved. That is, when the ratio of Ti and N (Ti / N) is in the range of 1.2 to 2.5, the amount of solid solution Ti is extremely reduced and the high temperature stability of the TiN precipitate is increased, so that the fine TiN precipitate of 0.01-0.1 μm or less is 0.5 μm or less. A result of distribution of 1.0 × 10 ⁷ pieces / mm 2 or more at intervals was obtained. This increases the content of nitrogen at the same Ti content, so that all the Ti atoms dissolved are easily combined with the nitrogen atom, and in a high nitrogen environment, the amount of solid solution Ti decreases, making TiN precipitates more stable at higher temperatures than in the case of low nitrogen content. It is analyzed that this is because the solubility product is lowered. In the present invention, the ratio of N / B, Al / N, V / N, and N and Ti + Al + B + (V) is considered, in view of the fact that it may promote aging due to the presence of solid solution N in a high nitrogen environment. Is managed as a whole to precipitate N into BN, AlN, and VN.

In addition, in the present invention, TiO is managed at 4-10 to precipitate and distribute TiO oxides that are stable even at 1400 ° C. or higher, thereby suppressing the growth of old austenite grains.

[2] ferrite grain size management

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

[3] microstructure of weld heat affected zone

In the present invention, the Ti / O ratio is controlled to 4-10 to appropriately proportion the ratio of Ti contained in the oxide, and a large amount of fine ferrite is produced in the former austenite mouth using precipitates such as AlN and BN. Ferrites are mostly polygonal ferrites and acicular ferrites, which greatly improve the toughness of heat-affected zones.

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

[Welding Structural Steels]

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

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

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

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

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

Mn deoxidizes in steel, improves weldability, hot workability and strength, and forms a solid solution to form a solid solution in the matrix to strengthen the matrix to secure strength and toughness. Do. However, when the Mn content is more than 2.0%, it has a detrimental effect on the toughness of the weld heat affected zone due to tissue heterogeneity due to manganese segregation rather than a solid solution strengthening effect. In addition, when the steel solidifies, macro segregation and micro segregation occur depending on the segregation mechanism, which promotes the formation of a central segregation zone in the center of rolling, thereby causing a low temperature transformation structure in the center of the base metal. In particular, manganese is an element that precipitates in the form of MnS around the Ti-based oxide and affects the formation of acicular and polygonal ferrites effective for improving the toughness of the weld heat affected zone.

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

Al is an element necessary as a deoxidizer and forms an Al oxide by reacting with oxygen to prevent Ti from reacting with oxygen, which not only helps Ti to form fine TiN precipitates, but also is effective for forming fine AlN precipitates in steel. . In order to form fine AlN precipitates, the Al content is preferably made 0.0005% or more. However, if it exceeds 0.1%, AlN is precipitated and the remaining solid Al promotes the formation of Widmanstatten ferrite and phase martensite, which are vulnerable to toughness during the cooling process of the weld heat affected zone. Lowers.

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

Ti is indispensable in the present invention because it combines with N to form a fine TiN precipitate that is stable at high temperatures. It is preferable to add more than 0.005% of Ti in order to obtain such a fine TiN precipitation effect, but when it exceeds 0.2%, coarse TiN precipitates and Ti oxides are formed in molten steel, and thus it is impossible to suppress the growth of the austenite grains of the weld heat affected zone. Because it is not desirable.

The content of boron (boron, B) is preferably limited to 0.0003-0.01%.

B is a very effective element for producing polygonal ferrite at grain boundaries as well as acicular ferrite having excellent toughness in grains. B forms a BN precipitate, which hinders the growth of the old austenite grains and forms Fe carbide in the grain boundary and in the mouth to promote ferrite transformation of acicular and polygons having excellent toughness. If the B content is less than 0.0003%, such an effect cannot be expected, and if it exceeds 0.01%, the hardenability increases, which may cause hardening of the weld heat affected zone and low temperature cracking.

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

N is an indispensable element for forming TiN, AlN, BN, VN, NbN, etc., and it is possible to minimize the growth of the old austenite grains in the weld heat affected zone during the high heat input welding and to increase TiN, AlN, BN, VN, NbN, etc. Increase the amount of precipitate; 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 acicular and polygonal ferrites, which are effective for improving the toughness of the weld heat affected zone. 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 indispensable for forming Ti oxide by reacting with Ti in molten steel. Ti oxide promotes the transformation of acicular ferrite during ferrite transformation from guustenite in the weld heat affected zone. If the O content is less than 0.0020%, it is not preferable because the number of Ti oxides is small, and if it exceeds 0.03%, coarse Ti oxides and other oxides such as FeO are produced, which is not preferable.

The ratio of Ti / N is preferably 1.2 to 2.5.

In the present invention, the Ti / N ratio is lowered to 2.5 or less, which has two advantages. First, it is possible to increase the amount of TiN, that is, the number of TiN precipitates. In other words, if the nitrogen content is increased at the same Ti content, all the Ti atoms dissolved in the cooling process during the playing process combine with the nitrogen atom, thereby increasing the fine TiN precipitation. Second, TiN is stable at high temperatures. That is, since the solubility product which shows the stability of the precipitate at high temperature such as the weld heat affected zone becomes smaller, the precipitate such as high nitrogen TiN is more stable than the case where the nitrogen content is low. On the other hand, if the Ti / N ratio is higher than 2.5, coarse TiN is crystallized in molten steel, which is a steelmaking process, and a uniform distribution of TiN is not obtained. Also, excess Ti remaining without precipitation as TiN remains in a solid solution to weld heat. Affects bad toughness. If the Ti / N ratio is less than 1.2, the amount of solid solution nitrogen in the base metal increases, which is detrimental to the toughness of the weld heat-oriented part.

It is preferable to make ratio of N / B into 10-40.

In the present invention, if the N / B ratio is less than 10, the precipitation amount of BN that promotes the ferrite transformation of polygons at the old austenite grain boundary during the post-weld cooling process is insufficient, and when the N / B ratio is over 40, the effect is saturated and dissolved. This is because the amount of nitrogen is increased to lower the toughness of the weld heat affected zone.

It is preferable to make Al / N ratio into 2.5-7.

In the present invention, when the Al / N ratio is less than 2.5, AlN precipitates for inducing needle-like ferrite transformation are insufficient, and the amount of solid solution nitrogen in the weld heat affected zone may increase, resulting in a weld crack, and an Al / N ratio of more than 7 In the case the effect is saturated.

It is preferable that ratio of (Ti + 2Al + 4B) / N is 6.5-14.

In the present invention, when the ratio of (Ti + 2Al + 4B) / N is less than 6.5, the growth inhibition of the austenite grain growth of the weld heat affected zone, the generation of fine polygonal ferrite at the grain boundary, the amount of solid solution nitrogen, the needle shape and the polygonal shape in the grain boundary Insufficient size and number of distribution of TiN, AlN, BN, and VN precipitates for ferrite formation and control of tissue fraction, and the effect is saturated when (Ti + 2Al + 4B) / N is more than 14 seconds. If V is added, the ratio of (Ti + 2Al + 4B + V) / N is preferably 7-17.

The ratio of Ti / O is preferably set to 4-10.

If the Ti / O ratio is less than 4, the number of TiO oxides required to inhibit austenite grain growth is insufficient, and the Ti ratio in TiO oxides becomes small, thus losing the function of acicular ferrite nucleation sites and improving the toughness of the weld heat affected zone. The effective acicular ferrite phase fraction is lowered. If the ratio of Ti / O is more than 10, the effect of inhibiting the weld heat affected zone austenite grain size crystallization is saturated, and the ratio of components such as Mn contained in the oxide becomes rather small, thus losing the function of the nucleation site of the ferrite in the mouth.

In order to further improve the toughness of the steel material (base material) and the heat-affected portion formed as described above, V is further added.

The content of vanadium (V) is preferably limited to 0.01-0.2%.

V is an element that combines with N to form VN to promote ferrite formation in the weld heat affected zone, and VN precipitates alone or precipitates in TIN precipitates to promote ferrite transformation. In addition, V combines with C to form VC, which acts to inhibit ferrite grain growth after ferrite transformation. When the V content is less than 0.01%, it is difficult to obtain the ferrite transformation promoting effect in the weld heat affected zone because the VN deposition amount is small. On the other hand, exceeding 0.2% is not preferable because it causes toughness of the base metal and the weld heat affected zone (HAZ) and improves the weld hardenability, which may cause the low temperature crack of the weld.

Moreover, it is preferable to make ratio of V / N into 0.3-9. In the present invention, when the V / N ratio is less than 0.3, it is difficult to secure an appropriate number and size of VN precipitates deposited and distributed on TiN + MnS precipitates for improving the toughness of the weld heat affected zone. If the V / N ratio exceeds 9, the size of the VN precipitates deposited at the TiN + MnS precipitate boundary is coarsened, and thus the VN precipitation frequency deposited at the TiN + MnS complex precipitate boundary is reduced, which is effective for the toughness of the weld heat affected zone. Reduce ferrite phase percentage.

In the present invention, in order to further improve the mechanical properties in the steel composition as described above, one or more selected from the group of Ni, Cu, Nb, Mo, Cr is further added.

The content of nickel (Ni) is preferably limited to 0.1-3.0%.

Ni is an effective element which improves the strength and toughness of the base material by solid solution strengthening. In order to achieve this effect, the Ni content is preferably 0.1% or more, but when the content exceeds 3.0%, the hardenability is increased to reduce the toughness of the weld heat affected zone and the possibility of high temperature cracking in the weld heat affected zone and the weld metal. This is not desirable because there is.

The content of copper (Cu) is preferably limited to 0.1-1.5%.

Cu is an element which is effective to secure the base material strength and toughness due to solid solution at the base. For this purpose, Cu content should be contained more than 0.1%, but if it exceeds 1.5%, it is not preferable because it increases the hardenability by increasing the hardenability in the weld heat affected zone and promotes high temperature crack in the weld heat affected zone and the weld metal. . In particular, Cu is an element that affects the formation of acicular and polygonal ferrites, which are effective in improving the toughness of the welded heat affected zone by depositing CuS around Ti-based oxides with sulfur, so that the content is 0.3-1.5%. desirable.

In addition, in the case of complex addition of Cu and Ni, the total sum thereof is preferably less than 3.5%. The reason is that less than 3.5% of the hardenability increases, which adversely affects the weld heat affected zone toughness and weldability.

The content of niobium (Nb) is preferably limited to 0.01-0.10%.

Nb is an effective element from the viewpoint of securing the strength of the base material. For this purpose, Nb is added in an amount of 0.01% or more. However, Nb is undesirable because it causes coarse precipitation of coarse NbC, which is detrimental to the toughness of the base material.

Chromium (Cr) is preferably made 0.05 to 1.0%.

Cr increases the hardenability and also improves the strength. If the content is less than 0.05%, the strength cannot be obtained and when the content exceeds 1.0%, the base metal and the HAZ toughness deteriorate.

Molybdenum (Mo) is preferably 0.05-1.0%.

Mo is also an element that increases the hardenability and improves the strength. The content thereof is 0.05% or more for securing the strength, but the upper limit is set to 1.0% like Cr for suppressing the HAZ hardening and the welding low temperature crack.

In addition, in the present invention, one or two kinds of Ca and REM are further added to suppress the grain growth of the austenite at the time of heating.

Ca and REM form an oxide having excellent high temperature stability, thereby suppressing the growth of the austenite grains when heated in the base metal and improving the toughness of the weld heat affected zone. In addition, Ca has the effect of controlling the coarse MnS shape during steelmaking. To this end, it is preferable to add more than 0.0005% of calcium (Ca) and more than 0.005% of Rem, but when Ca exceeds 0.005% of Rem of more than 0.05%, large inclusions and clusters are generated to harm the cleanliness of the steel. As REM, 1 type, or 2 or more types, such as Ce, La, Y, and Hf, may be used, and any of the above effects can be obtained.

· Microstructure of steel

In the present invention, after the hot rolling, the microstructure of the steel is ferrite + pearlite, and the size of the ferrite grains is preferably 20 μm or less. The reason is that when the grain size of the ferrite is larger than 20 μm, the size of the old austenite grains in the weld heat affected zone becomes 80 μm or more during high heat input welding, which is detrimental to the weld heat affected zone toughness.

In addition, the higher the percentage of ferrite in the composite structure of ferrite + perlite, the higher the toughness and elongation of the base metal, and the ferrite is 20% or more and most preferably 70% or more.

Precipitate

The former austenite grains of the weld heat affected zone are greatly influenced by the size, number and distribution of oxides or nitrides distributed in the base material when the austenite grain size of the base material is constant. In addition, since 30-40% of the nitrides distributed in the base material are welded to the base material at the time of high heat input welding (above the heating temperature of 1400 ℃ or more), the effect of inhibiting the growth of the austenite grains in the weld heat affected zone is reduced. There is a need for a uniform distribution of further nitrides taking into account the re-used nitrides. In order to suppress the growth of the old austenite in the weld heat affected zone, it is important to uniformly distribute the fine TiN precipitate to suppress the Ostwald ripening phenomenon in which some precipitates are coarsened. For this purpose, the TiN precipitates should be controlled to 0.5 μm or less to make the TiN distribution uniform.

In addition, it is preferable to limit the particle diameter and the critical number of TiN to 0.01-0.1 μm and 1.0 × 10 ⁷ or more per 1 mm ² . The reason for this is that less than 0.01 μm is easily re-used to the base metal during high heat input welding, and the effect of inhibiting the growth of the old austenite grains is insufficient. When the thickness exceeds 0.1 μm, pinning of the old austenite grains This is because the growth inhibition effect is reduced and behaves like coarse non-metallic inclusions, which adversely affects mechanical properties. In addition, when the number of precipitates is less than 1.0 × 10 ⁷ per 1 mm ² , it is difficult to control the size of the old austenite grains of the weld heat affected zone at the time of welding higher than the heat input to be 80 μm or less, which is a threshold value.

·oxide

According to the present invention, it is preferable to control the ratio of Ti / O to 4-10 so that the particle size and the critical number of the Ti oxide in the base material are in the range of 0.5-2.0 μm and 1.0 × ^{10 2} -1.0x10 ³ per mm ² . The reason for this is that when the oxide size is less than 0.5 µm, the needle ferrite nucleation promoting effect is insufficient in the weld heat-affected grains. This is because the amount of transformation of the acicular ferrite is reduced. If the number of precipitates is 1.0X10 ² or less per 1 mm ^2, it is not preferable because the amount of acicular ferrite formed from Ti oxide is insufficient. If the amount of precipitate is more than 1.0X10 ^3, acicular ferrite formed from Ti oxide is formed from adjacent Ti oxide. It is not preferable because it overlaps with acicular ferrite and interferes with formation of acicular ferrite.

[Method of manufacturing welded structural steel]

Refining Process (Deoxidation 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-of-furnace refining, degassing (RH process) is performed. Normal deoxidation takes place in primary and secondary refining.

In the present invention, the dissolved oxygen is adjusted to an appropriate level in such a deoxidation process and then Ti is added to obtain a distribution of Ti oxide at an appropriate level. Based on the fact that the amount of dissolved oxygen greatly affects the production behavior of the oxide, the present inventors have found that there is an adequate amount of dissolved oxygen for producing a large number of oxides as a result of obtaining a desired Ti oxide. The value was confirmed to be 50-200ppm. In order to obtain this level of dissolved oxygen, a deoxidizer, which has a stronger deoxidizing power than Ti, should be added before the introduction of Ti to prevent the formation of coarse Ti oxide in molten steel while increasing the number of Ti oxides during the play. These Ti oxides can increase the precipitation frequency of precipitates, such as TiN, MnS, CuS, VN.

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 Ti, a deoxidizer with a greater deoxidizing power than Ti must be added to ensure proper dissolved oxygen (50-200 ppm) to shorten the progress of the deoxidation process, and to secure proper dissolved oxygen to the desired TiO oxide. While forming Ti can be made into a TiN precipitate. That is, since the deoxidizer having a higher deoxidizing power than Ti is pre-injected, the ratio of Ti remaining in solid solution becomes considerably higher than that of coarse primary oxides in molten steel, so that a fine amount of Ti oxide and TiN nitride during solidification Will be able to secure.

In the present invention, before adding an element (Al, REM, Zr, Ca, Mg) having a higher deoxidizing power than Ti, deoxidation 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. This method has the advantage of easy control of the appropriate dissolved oxygen amount. As an example of the deoxidation operation pattern in the present invention, Mn, Si, Al in the first refining and then injecting Al in the second refining and Ti is added, this addition of Ti is vacuum degassing in the second refining In the case of adopting a gas treatment step, it is also possible to inject it after vacuum degassing.

In the present invention, it is preferable to add a deoxidation element having a greater deoxidizing power than Ti before deoxidizing the dissolved oxygen amount to 50-200 ppm. The reason is that when the dissolved oxygen amount is less than 50 ppm, the amount of oxygen in the molten steel is too small to form an adequate amount of Ti-based oxide for showing the effect of the present invention. In addition, it is difficult to obtain an alloying system composed of accurate chemical components by oxidizing other alloying elements.

Subsequently, it is preferable to add Ti so that the content is 0.005-0.2%, which improves the toughness of the weld heat affected zone by promoting the acicular ferrite transformation as well as the proper distribution of Ti oxide and Ti precipitates for showing the effect of the present invention. It is the range which considered the effective precipitate behavior of MnS, CuS, BN, VN, etc. together.

As described above, after the addition of Ti, the molten steel is refined, at which time the refining is performed for 0.5-20 minutes. When the retention time of the molten steel is less than 0.5 minutes, the coarse Ti-based primary oxide does not have time to separate the flotation and remains after coagulation, and the average size thereof increases, and when it exceeds 20 minutes, the oxides are bonded to each other. Therefore, it is difficult to secure an appropriate amount of solid solution Ti in the molten steel to coarsen and exhibit the effects of the present invention.

Casting process

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

According to the study of the present invention, the continuous casting speed is less than 1.1 m / min, more preferably about 0.9 to 1.1 m / min, which is lower than the usual casting speed of about 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 the productivity is lowered. When it is faster than 1.1m / min, cast surface cracks are more likely to occur.

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

Hot rolling process

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

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

After rolling as described above, in order to prevent ferrite grain growth, it is cooled at a rate of 1 ° C./min or more up to a ferrite transformation end temperature of ± 10 ° C. Preferably, it is cooled at a rate of 1 ° C./min or more to the ferrite transformation end temperature and thereafter air-cooled. Of course, it is possible to cool the ferrite transformation temperature below the end temperature, that is, above 1 ° C / min to room temperature, but it is uneconomical. If the cooling rate is less than 1 ° C / min, it leads to grain growth of the recrystallized fine ferrite, it is difficult to secure the ferrite grain size of the base material 20㎛ or less.

[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 have a size of 0.01-0.1 μm and 1.0 × 10 ⁷ or more per 1 mm ² , and the interval is 0.5 μm or less, and TiO oxides are finely distributed.

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 must be controlled to be 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 20 μm or less has a phase fraction of 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 an Al-based composite oxide according to the present invention.

Hereinafter, the present invention will be described in detail by way of examples.

EXAMPLE

Steel grades having the composition as shown in Table 1 were prepared as slabs by a continuous casting method by dissolving them in a converter, and the composition ratio between the alloying elements for each steel type to show the effect of the present invention is shown in Table 2. Table 3 shows the solidification rate of slab, slab heating temperature, heating time, rolling start temperature and end temperature, rolling reduction, and rolling rate of 25 ~ 40mm thickness in rolling process. At this time, the rolling reduction ratio of all steel grades was made 75% or more.

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

Tensile test piece was used KS standard (KS B 0801) No. 4 test piece and the tensile test was tested at the cross head speed (5mm / mim). The impact test piece was manufactured according to KS (KS B 0809) No. 3 test piece, and the notch direction was processed on the side of the rolling direction (L-T) in the case of the base material and in the welding line direction on the welding material. In addition, in order to investigate the austenite grain size according to the maximum heating temperature of the welding heat affected zone, it is heated to 140 ℃ / sec condition for 1 second after the heating up to the maximum heating temperature (1200 ~ 1400 ℃) by using the simulation welding simulator (simulator). After holding, it was quenched with He gas. The quenched specimens were ground and corroded to determine the austenite grain size at the highest heating temperature condition by KS (KS D 0205).

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

Impact toughness evaluation of the weld heat affected zone is 800-500 ℃ cooling after heating the welding conditions equivalent 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.

Steel grade used division Primary deoxidation sequence Molten oxygen (ppm) Ti addition amount (%) Molten steel holding time (min) Oxygen content after coagulation (ppm) Casting condition Casting speed (m / min) Specific quantity (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 58 0.015 15 30 1.0 0.35 Invention 2 Mn → Si 55 0.015 15 34 1.0 0.35 Invention 3 Mn → Si 65 0.015 15 32 1.0 0.35 Comparative Material 1 Mn → Si 26 0.015 15 8 1.0 0.35 Comparative Material 2 Mn → Si 206 0.015 65 124 1.0 0.35 Inventive Steel 2 Invention 4 Mn → Si 69 0.053 16 54 0.95 0.32 Invention Steel 3 Invention 5 Mn → Si 58 0.017 15 32 1.0 0.32 Inventive Steel 4 Invention 6 Mn → Si 54 0.022 14 31 1.0 0.32 Inventive Steel 5 Invention 7 Mn → Si 69 0.053 15 52 1.1 0.35 Inventive Steel 6 Invention Material 8 Mn → Si 55 0.022 15 32 1.0 0.35 Inventive Steel 7 Invention Material 9 Mn → Si 61 0.016 16 36 1.1 0.35 Inventive Steel 8 Invention 10 Mn → Si 72 0.019 18 44 1.0 0.35 Inventive Steel 9 Invention 11 Mn → Si 67 0.021 14 43 0.98 0.32 Inventive Steel 10 Invention Material12 Mn → Si 75 0.026 11 52 1.0 0.32 Inventive Steel 11 Invention Material 13 Mn → Si 68 0.020 12 48 1.1 0.35 Conventional steel (1-11) is not specifically described its manufacturing conditions

Steel grade used division Heating temperature (℃) Heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount / accumulated loading amount at recrystallization area (%) Cooling rate (℃ / min) Cooling end temperature (℃) Invention 2 Inventive Example 1 1150 170 1030 780 65/80 7 450 Inventive Example 2 1200 130 1040 790 65/80 7 450 Inventive Example 3 1240 90 1040 780 65/80 7 450 Comparative Example 1 1050 60 1040 780 65/80 7 450 Comparative Example 2 1300 250 1035 780 65/80 7 450 Invention 1 Inventive Example 4 1200 130 1020 790 65/80 6 450 Invention 3 Inventive Example 5 1200 130 1040 790 65/80 6 450 Comparative Material 1 Comparative Example 3 1210 120 1030 780 65/80 0.1 Room temperature Comparative Material 2 Comparative Example 4 1210 120 1030 790 65/80 19 Room temperature Invention 4 Inventive Example 6 1180 150 1020 790 60/80 7 450 Invention 5 Inventive Example 7 1190 140 1010 800 60/80 8 450 Invention 6 Inventive Example 8 1220 110 1010 810 60/75 7 450 Invention 7 Inventive Example 9 1220 110 1020 800 60/75 10 450 Invention Material 8 Inventive Example 10 1210 120 1010 790 60/75 10 450 Invention Material 9 Inventive Example 11 1220 110 1000 780 55/70 10 450 Invention 10 Inventive Example 12 1210 120 1010 790 55/70 9 450 Invention 11 Inventive Example 13 1230 100 1000 800 55/70 8 450 Invention Material12 Inventive Example 14 1220 110 1020 780 55/70 10 450 Invention Material 13 Inventive Example 15 1210 130 1020 780 65/75 10 450 Conventional Steel 11 1200 - Ar ₃ or higher 960 80 Cooling Manufacturing conditions of conventional steel (1-10) are not specifically presented

As shown in Table 5, the number of precipitates (Ti-based nitride) of the hot rolled material produced by the present invention has a range of 2.5 × ^{10 8} / mm ² or more, whereas in the case of conventional steel 4.07 X10 ⁶ / mm ² Since the following ranges show that the invention material has a fairly uniform and fine precipitate size compared to the conventional material, the number is also significantly increased. In addition, the present invention has a number of oxides (TiO) of about 1.7x10 ² ~ 3.26x10 ³ ranges and the average size is about 1.2-1.8㎛ range. On the other hand, in the structure of the base steel of the present invention, the steel of the present invention has a ferrite grain size (FGS) of about 9-12 μm, which is very fine compared to that of the comparative material. It consists of fractions.

Table 6 shows the properties of the weld heat affected zone of the present invention steel and conventional steel. The austenitic grain size at the maximum heating temperature of 1400 ° C. such as the welding heat affected zone shows that the present invention has a range of 52-64 μm, while the conventional material has a very coarse range of about 180 μm or more. Can be. Therefore, in the present invention, it can be seen that the austenite grain suppression effect of the weld heat affected zone during welding is very excellent. In addition, the ferrite phase fraction of the steel of the present invention at a heat input of 100 kJ / cm is about 70% or more.

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

Claims (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%, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤ 7, 6.5 ≤ (Ti + 2 Al + 4B) / N ≤ 14, 4 ≤ Ti / O ≤ 10, is composed of the remaining Fe and other impurities, the microstructure of the ferrite and pearlite composite structure of 20㎛ or less Welded structural steel with excellent weld heat affected zone toughness. 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. Welded structural steels with excellent toughness of weld heat affected zone, characterized in that the satisfactory. 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%. A welded structural steel with excellent toughness in welding heat affected zone, characterized by containing one or two or more selected from the group selected from Ca: 0.0005-0.005% and REM: 0.005 to 0.05%. The TiN precipitate of 0.01-0.1 μm is distributed in the steel material at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less, and the TiO oxide of 0.5-2.0 μm is 1 × 10 ² -1 × 10 ^3. Welded structural steels with excellent toughness of weld heat affected zones, characterized by distribution of pieces / mm ² . 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%, 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, and after heating the steel slab with the remaining Fe and other impurities in the range of 1100-1250 ℃ for 60-180 minutes A method for manufacturing structural steel with excellent toughness in weld heat affected zone, comprising hot rolling at a recrystallization zone of 40% or more, followed by cooling to a ferrite transformation end temperature of ± 10 ° C 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 with excellent weld heat affected zone toughness 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%. 1 or 2 or more selected from the group of Ca: 0.0005-0.005%, Rem: 0.005 to 0.05%, the method for producing a welded structural steel having excellent toughness of the weld heat affected zone, characterized in that it contains. . The steel slab according to claim 5 or 6, wherein the deoxidation element having a greater deoxidizing power than Ti is added to molten steel before Ti is introduced, and the dissolved oxygen content is deoxidized to 50 to 200 ppm, and Ti is added so that Ti is 0.005 to 0.2%. And degassing and refining the continuous casting process. The method of claim 5 or 6, wherein the continuous casting of the slab is molten steel A method for producing a welded structural steel with excellent toughness in welding heat affected zones, characterized by mild cooling with a specific amount of 0.3 to 0.35 l / kg in a secondary cooling zone while casting at a speed of 0.9 to 1.1 m / min. The welding is applied to the steel (base material) of claim 4 to form a prior austenite of 80 μm or less in the weld heat affected zone, and then quenched to 70% of the ferrite of 20 μm or less in the microstructure of the weld heat affected zone. Welding structure with excellent toughness of the weld heat affected zone composed of the above ordinary fraction. The weld structure of claim 10, wherein the ferrite is polygonal ferrite and acicular ferrite.