KR100979007B1 - Ultra-High Strength Steel Sheet For Line Pipe Having Excellent Low Temperature Toughness And Method For Manufacturing The Same - Google Patents

Abstract

The present invention relates to a steel sheet for ultra-high strength line pipe having excellent resistance to unstable ductile fracture and brittle fracture propagation at cryogenic temperatures and a method for manufacturing the same. More specifically, DBTT (ductile-brittle transition temperature) of ultra-high strength steel by Ni addition It is an object of the present invention to provide a steel sheet having a tensile strength of 930 MPa or more and excellent Charpy impact toughness and DWTT characteristics, and a method of manufacturing the same, through precise management for obtaining high strength-toughness structure in a reduction, rolling, and cooling process.

Steel sheet of the present invention in weight%, C: 0.03-0.10%, Si: 0.05-0.6%, Mn: 1.6-2.1%, Cu: 0.5% or less (not containing 0%), Ni: 0.8-2.0% , Nb: 0.02 to 0.06%, V: 0.1% or less (without 0%), Mo: 0.2-0.5%, Cr: 1.0% or less (without 0%), Ti: 0.005 to 0.03%, Al: 0.01% to 0.05%, B: 0.0005% to 0.0020%, N: 0.002% to 0.006%, Ca: 0.0005% to 0.006%, P: 0.02% or less (not including 0%), S: 0.005% or less (0% Do not include), the remainder is characterized by containing Fe and other unavoidable impurities, the tensile strength of 930MPa or more and the impact toughness of 200J or more. Moreover, the manufacturing method of the said steel plate is characterized by the above-mentioned.

Line pipe, cryogenic toughness, tensile strength, impact toughness, DWTT characteristics

Description

Ultra-High Strength Steel Sheet For Line Pipe Having Excellent Low Temperature Toughness And Method For Manufacturing The Same}

The present invention relates to a steel sheet for ultra-high strength line pipes having excellent resistance to unstable ductile fracture and brittle fracture propagation at cryogenic temperatures, and a method of manufacturing the same. More specifically, precisely controlling alloy components and rolling methods to improve strength and cryogenic toughness. This invention relates to a line pipe steel sheet having a strength of 930 MPa or more and excellent in Charpy impact toughness and DWTT characteristics and a method of manufacturing the same.

The line pipe means a steel pipe buried in the ground mainly for transportation of crude oil or natural gas, and high pressure acts on the line pipe because high pressure gas or crude oil flows in the line pipe. In addition, in order to increase the efficiency of the line pipe, it is necessary to increase the amount of crude oil or gas (hereinafter, referred to as 'crude oil') that can be transported per unit time. There is a need. Until now, it is a fact that steel sheets of X70 and X80 grades are mainly used in view of the strength standards of line pipes. The steel sheet of grade X70 has a yield strength of about 70 ksi or about 480 MPa, and the steel sheet of grade X80 has a yield strength of about 80 ksi or about 550 MPa. However, the use of these grades of steel is not economical as long distance pipeline projects will inevitably require an increase in sheet thickness.

In addition, low-temperature toughness requirements for steel line pipes have been strengthened by the development of cold-field oil fields, such as the development of oil fields in Siberia and Alaska, which have poor weather conditions. Demanding steel. In order for line pipe steels to be used safely at low temperatures, the DWTT characteristics, which represent Charpy impact toughness and brittle fracture stopping characteristics, which are unstable ductile fracture resistance must be excellent. In the conventional environment, it was possible to use Charpy impact toughness at -40 ℃ to 160J or more and DWTT ductile fracture rate at -20 ℃ to 85% or more. However, Charpy impact toughness at -60 ℃ is very high in cold environment. More than 200J, steel materials with a DWTT ductile fracture rate of more than 85% at -40 ° C are required.

Accordingly, in recent years, steel fabrication technology for ultra-high strength line pipe of X100 (yield strength 100ksi or more) or X120 (yield strength 120ksi or more) has been dramatically improved. As a representative technique, by adding B together with Mo, the hardenability is remarkably improved, and the TMCP (mechanical method for changing the physical properties of steel sheet to desired properties by applying mechanical history to steel sheet at the same time) is applied. X120 ultra-high strength line pipe steel manufacturing method. However, these techniques are due to the deterioration of toughness due to high strength. Charpy impact toughness can be guaranteed at -40 ℃ and DWTT can be guaranteed at -20 ℃.

The inventors of the present invention, while considering the equipment and plate deformation, and easy to mass-produce, and looking for a way to manufacture a steel sheet for ultra-high strength line pipe with excellent cryogenic toughness, when Ni is added in ultra-high strength steel DBTT (ductile-brittle transition By dramatically lowering the temperature), it has been confirmed through experiments that the excellent impact toughness and DWTT characteristics are shown without deterioration in strength even at cryogenic temperatures, and the present invention has been proposed based on the experimental results.

The present invention is characterized in that the Charpy impact toughness and DWTT have a strength of 930 MPa or more through precise management for obtaining DBTT (ductile-brittle transition temperature) of ultra-high strength steel by Ni addition, high strength-toughness structure in rolling and cooling processes. It is an object of the present invention to provide a steel sheet having excellent characteristics and a method of manufacturing the same.

In order to achieve the above object, the steel sheet for line pipe of the present invention is a weight%, C: 0.03-0.10%, Si: 0.05-0.6%, Mn: 1.6-2.1%, Cu: 0.5% or less (not including 0%). Ni: 0.8 to 2.0%, Nb: 0.02 to 0.06%, V: 0.1% or less (does not contain 0%), Mo: 0.2 to 0.5%, Cr: 1.0% or less (does not contain 0%) ), Ti: 0.005 to 0.03%, Al: 0.01 to 0.05%, B: 0.0005 to 0.0020%, N: 0.002 to 0.006%, Ca: 0.0005 to 0.006%, P: 0.02% or less (not including 0%) , S: 0.005% or less (does not contain 0%), the rest is characterized by containing Fe and other unavoidable impurities, tensile strength of 930MPa or more and impact toughness of 200J or more.

Furthermore, the present invention reheats the steel slab having the composition range at 1050 ~ 1200 ℃, in the austenitic recrystallization zone, the average reduction rate per pass more than 10% once or two or more multi-stage rolling after the austenite recrystallization At the rate of 20 ~ 50 ℃ / sec after finishing rolling with 70 ~ 90% cumulative pressure drop in one or two or more multi-stages at a temperature lower than the temperature and higher than the temperature (Ar ₃ ) where austenite is converted into ferrite (Ar ₃ ) After cooling and ending the cooling at 200 ~ 400 ℃, it is characterized by air-cooling to room temperature.

According to the present invention, by controlling the rolling in the austenite recrystallization and unrecrystallization region, the effective particle size is extremely fine, and the addition of Ni promotes bainitic ferrite transformation, which is excellent in toughness, thereby securing ultra high strength and thus cryogenic toughness. It is possible to provide an excellent steel sheet for high strength line pipe.

Hereinafter, the composition range of this invention is demonstrated concretely.

C: 0.03 ~ 1.0 wt%

C is the most effective element for strengthening the matrix of metals and welds through solid solution strengthening, and it is possible to obtain strengthening effect by precipitation hardening by formation of small size cementite, V and Nb carbonitride and Mo carbide. In addition, Nb carbonitride can simultaneously improve strength and low temperature toughness through grain refinement by inhibiting austenite recrystallization and preventing grain growth during hot rolling. C also serves to improve the hardenability to form a strong microstructure inside the steel sheet during cooling. In general, when the amount is less than 0.03% by weight, such a strengthening effect cannot be obtained, and when added in excess of 0.1% by weight, the toughness of the base metal and the welded heat affected zone including the low temperature crack after the spot welding is reduced. More preferably, it is good to add 0.04-0.08 weight%.

Si: 0.05 ~ 0.6 wt%

Si plays a role of deoxidizing molten steel by assisting Al and also has an effect as a solid solution strengthening element. If the addition amount is less than 0.05%, the deoxidation of the molten steel may not be sufficient, so the toughness may be reduced. When the Si content is excessively added in excess of 0.6% by weight, the on-site weldability and the toughness of the weld heat affected zone are greatly reduced.

Mn: 1.6 ~ 2.1 wt%

Mn is an effective element to solidify the steel to be added at least 1.6% by weight can exhibit high strength with the effect of increasing the hardenability. However, the addition of more than 2.1% by weight promotes central segregation and lowers toughness when casting slabs in the steelmaking process. In addition, the addition of excessive Mn excessively improves the curing ability, worsens the field weldability, thereby lowering the toughness of the weld heat affected zone.

Cu: 0.5 wt% or less (does not contain 0%)

Cu is an element to strengthen the base metal and the weld heat affected zone. However, when Cu is added excessively, the toughness and spot weldability of a weld heat affected zone will fall. In addition, when Cu exceeds 0.5% by weight, hot-working cracks due to the Cu film at the grain boundary during rolling are generated, so it is limited to 0.5% or less.

Ni: 0.8-2.0 wt%

Ni is a key element in the present invention and is an element that improves physical properties without losing on-site weldability and low temperature toughness. Compared with Mn and Mo, Ni forms less hard phases such as island martensite, which lowers the low temperature toughness, and improves the toughness of the weld heat affected zone. In addition, as the Ni content increases, the content of bainitic ferrite, which is an effective tissue for cryogenic toughness, increases, thereby dramatically improving impact toughness and DWTT characteristics. If the Ni content is less than 0.8%, the bainitic ferrite fraction effective for cryogenic toughness is lowered, making it difficult to satisfy the toughness required at cryogenic temperatures. However, Ni is an expensive element and rather lowers the toughness of the weld heat affected zone when Ni is added in excess of 2.0% by weight. Therefore, the Ni content is 0.8 to 2.0% by weight.

Nb: 0.02 ~ 0.06 wt%

Nb plays a role of simultaneously improving strength and toughness through grain refinement. Nb carbonitride produced during hot rolling suppresses austenite recrystallization and prevents grain growth, thereby making fine austenite grains. When added together with Mo, the austenite recrystallization is suppressed to increase the grain refining effect, the reinforcing effect through the precipitation strengthening and the hardenability is more pronounced. When B is present, the effect of further increasing the hardenability can be obtained. In order to obtain such an effect, it should contain at least 0.02% by weight. However, when added in excess of 0.06% by weight, it is difficult to expect the effect increase any more, and also adversely affect the weldability and the weld heat affected zone toughness. More preferably, it is good to add 0.02-0.06 weight%.

V: 0.1 wt% or less (does not include 0%)

V plays a similar role to Nb, but the effect is somewhat weaker than Nb. However, the effect is greatly magnified when Nb and V are added together. However, considering the toughness and weldability of the weld heat affected zone, the upper limit thereof is made 0.1 wt%. More preferably, 0.08 wt% or less is added.

Mo: 0.2 ~ 0.5 wt%

Mo improves the hardenability, especially when added together with B, the effect of improving hardenability is very large. In addition, when added together with Nb suppresses austenite recrystallization contributes to grain refinement. However, excessive addition of Mo lowers the toughness of the weld heat affected zone during the field welding, so it should be kept below 0.5% by weight. If the Mo content is less than 0.2%, the bainitic ferrite content required for securing strength and toughness is lowered due to a decrease in hardenability, which is not preferable.

Cr: 1.0 wt% or less (does not include 0%)

Cr plays a role of improving the hardenability. However, the addition of excessive Cr causes low temperature cracks after welding in the field, which degrades the toughness of the base metal and the heat affected zone of the weld, so it should be maintained at 1.0 wt% or less.

Ti: 0.005 ~ 0.03 wt%

Ti contributes to grain refinement by forming fine Ti nitride (TiN) to suppress austenite grain coarsening during slab heating. In addition, TiN not only prevents grain coarsening in the weld heat affected zone, but also improves toughness by removing N in molten steel. In order to sufficiently remove N, Ti must be at least 3.4 times the amount of N added. Therefore, Ti is a very useful element to refine the strength and grains of the base metal and the weld heat affected zone, and is present as TiN in the steel to inhibit the growth of grains during heating for rolling. The solid solution in the steel combines with carbon to form precipitates of TiC, and the formation of TiC is very fine, greatly improving the strength of the steel. When the amount of Al added is very small, Ti oxide is formed to act as a nucleation site of intragranular acicular ferrite in the weld heat affected zone. Therefore, it is necessary to add at least 0.005% by weight or more in order to obtain the austenite grain growth inhibition effect by TiN precipitation and the strength increase by TiC formation. On the other hand, when added in excess of 0.03% by weight, the coarsening of Ti nitride and the hardening by Ti carbide are excessively harmful to low temperature toughness. Since the toughness of the affected portion deteriorates, the upper limit of Ti addition is made 0.03% by weight. More preferably 0.01 to 0.02% by weight is added.

Al: 0.01 ~ 0.05 wt%

Al is generally added for the purpose of deoxidation of the steel. Further, not only the microstructure is fine but also the toughness of the heat affected zone is improved by removing N from the coarse grain region of the weld heat affected zone. However, when contained in excess of 0.05% by weight, Al oxide (Al ₂ O ₃ ) is formed to reduce the toughness of the base metal and the heat affected zone. If it is less than 0.01%, since the deoxidation effect is inadequate and toughness falls, it is preferable to set it as 0.01 to 0.05%.

B: 0.0005 ~ 0.0020 wt%

B greatly improves the hardenability in low carbon steels and increases weldability and low temperature crack resistance. In particular, it serves to increase the effect of improving the hardenability of Mo and Nb and increases the strength of the grain boundary to suppress intragranular cracking generated by hydrogen. If the B content is less than 0.0005%, the bainitic ferrite content required for securing strength and toughness is lowered due to a decrease in hardenability. However, excessive addition of B causes embrittlement by Fe ₂₃ (C, B) ₆ precipitation. Therefore, the content of B should be determined in consideration of the content of other hardenable elements. In the present invention, the content of B is preferably in the range of 0.0005 to 0.0020% by weight as described above.

N: 0.002 ~ 0.006 wt%

The reason for component limitation of N is attributable to the above Ti addition. In general, N is dissolved in the steel and precipitates to increase the strength of the steel, which is much greater than carbon. However, on the other hand, it is known that toughness decreases as more nitrogen exists in steel, and it is a general trend to reduce nitrogen content as much as possible. However, in the present invention, it is not preferable to reduce N excessively because the proper amount of nitrogen is present to react with Ti to form TiN and to inhibit grain growth during reheating. If the N content is less than 0.002%, the content of TiN precipitates is small, so that the particle size growth inhibitory effect is greatly reduced, so the lower limit of the N content is 0.002%. If the N content is too large, N is present in solid solution N, not in TiN form, so that the toughness is greatly reduced. Therefore, the upper limit is made 0.006% or less.

Ca: 0.0005 ~ 0.006 wt%

Ca is mainly used as an element to control the shape of MnS inclusions and to improve low temperature toughness. However, the Ca addition exceeding 0.006% by weight of large amounts of CaO-CaS forms and combines to form coarse inclusions, thereby lowering the cleanliness of the steel and damaging the field weldability. If the amount of Ca is less than 0.0005%, the shape control of the MnS inclusions is insufficient, so the lower limit of Ca is set to 0.0005%.

P: 0.02 wt% or less (does not include 0%)

P is an impurity element present in steel, and segregation at the center of the steel sheet mainly damages toughness. Therefore, it is necessary to actively reduce P. However, in order to reduce P to an extreme, the steelmaking process load is intensified and the problem is less than 0.02 wt%. Since it does not generate | occur | produce greatly, the upper limit is made into 0.02 weight%.

S: 0.005 wt% or less (does not include 0%)

S combines with Mn to form a non-metallic inclusion to embrittle steel and causes red-brittle brittleness. Like S, P is limited to an upper limit of 0.005% by weight in consideration of steelmaking process load. More preferably, 0.002 wt% or less is added.

Other than the above components may include Fe and other unavoidable impurities.

Hereinafter, the manufacturing method of this invention is demonstrated concretely.

In the manufacturing method of the present invention, the slab is heated, and then the heated slab is subjected to one or two or more multi-stage rollings in an austenite recrystallization zone with an average reduction ratio of 10% or more per pass, and then the austenite recrystallization temperature. After finishing rolling at 70 ~ 90% of the cumulative reduction ratio in one or two or more stages at a temperature higher than the temperature at which austenite is transformed into ferrite (Ar ₃ ), it is cooled at a speed of 20-50 ° C./sec. It consists of the process of terminating cooling at 200 ~ 400 ℃. It is preferable to air-cool or to cool a steel plate below the said cooling end temperature.

Hereinafter, detailed conditions of each step will be described.

(1) Slab heating: 1050 ~ 1200 ℃

The heating process of the slab is a process of smoothly performing the subsequent rolling process and heating the steel to sufficiently obtain the properties of the target steel sheet, so the heating process should be performed within an appropriate temperature range according to the purpose. What is important in the heating process is that not only the precipitated elements inside the steel sheet should be heated uniformly so as to be sufficiently dissolved, but also the maximum protection against excessive coarsening of grains due to too high a heating temperature is important. If the heating temperature of the steel is less than 1050 ° C., Nb or V may not be re-used in the steel, making it difficult to achieve high strength of the steel sheet, and partial recrystallization may occur, causing austenite grains to be unevenly formed, making it difficult to toughen. When the temperature exceeds 1200 ° C., the austenite grains are excessively coarse, thereby providing a cause of an increase in grain size of the steel sheet, and as a result, the toughness of the steel sheet is extremely deteriorated. Therefore, it is preferable that the suitable heating temperature range is 1050-1200 degreeC.

(2) rolling conditions

In order for the steel sheet to have low temperature toughness, the austenite grains must be present in a fine size, which is possible by controlling the rolling temperature and the reduction ratio. In the present invention, the rolling is preferably carried out in two temperature ranges, the recrystallization behavior is different in the two temperature ranges, it is preferable to set the conditions respectively. First, 20-60% of the initial slab thickness is subjected to one rolling or two or more multistep rolling in the austenite recrystallization region. Rolling in the austenitic recrystallization area as described above has the effect of reducing the grain size through austenite recrystallization, in the case of multi-stage rolling to reduce the reduction rate and time of each step so that grain growth does not occur after austenite recrystallization Control. The degree of austenite recrystallization is governed by the reduction ratio and temperature in the austenite recrystallization region. At this time, if the reduction ratio of each step is too low, recrystallization only partially occurs, so that austenite grains are not uniform and satisfactory crack growth resistance cannot be obtained. Therefore, the reduction ratio of each step is limited to 5% or more. In addition, as the rolling progresses and the temperature decreases, the reduction ratio must be increased, and the average reduction ratio (the total of each reduction ratio divided by the reduction stage number) in the austenite recrystallization zone must be 10% or more. On the other hand, when the holding time between each pressing step is excessively long, grain growth occurs and uniformity and miniaturization of austenite grain size cannot be obtained and excellent crack growth resistance cannot be obtained. Thus, grain growth occurs after austenite recrystallization. To avoid this, the holding time between each pressing step should be within 20 seconds. Fine austenite grains formed by the above process serves to improve the crack growth resistance of the final plate.

Thereafter, one rolling or two or more multistage rollings are performed in the austenite unrecrystallized region between T _nr (temperature at which austenite recrystallization does not occur) and Ar ₃ temperature (temperature at which austenite is converted to ferrite). At this time, rolling is performed at 70% or more of the slab thickness after rolling in the recrystallization temperature range. As the amount of reduction in the unrecrystallized region increases, rolling is more favorable, but rolling in excess of 90% is not possible due to the temperature drop. Therefore, the upper limit of the reduction ratio is 90%. This rolling between T _nr (temperature at which austenite recrystallization does not occur) and Ar ₃ temperature (temperature at which austenite is transformed into ferrite) distorts the grains and develops dislocations due to deformation in the grains. It acts as a nucleation site that forms low temperature transformation phase

(3) Cooling rate: 20 ~ 50 ℃ / sec

Cooling rate is an important factor to improve the toughness and strength of the steel sheet. The cooling rate is to control the structure of the steel sheet as bainitic ferrite or acicular ferrite as described above, when the cooling rate is less than 20 ℃ / sec undesired tissue such as Polygonal Ferrite (Polygonal Ferrite) It is formed while having a coarse grain size, and there is a fear that the strength and toughness are greatly reduced to a level not meeting the target value of the present invention. On the contrary, when cooling at a high cooling rate exceeding 50 ° C./sec, the warpage of the steel sheet occurs due to the excessive amount of cooling water, resulting in poor shape of the steel sheet.

(4) Cooling end temperature: 200 ~ 400 ℃

In order to control the internal structure of the steel sheet, it is necessary to cool it to a temperature at which the effect of the cooling rate is sufficiently manifested. If the cooling stop temperature, which is the temperature at which the cooling stops, exceeds 400 ° C., it is difficult to sufficiently form bainitic ferrite and acicular ferrite having fine grains in the steel sheet, and thus the effect of improving the tensile strength is insufficient. Therefore, the upper limit of the cooling stop temperature needs to be limited to 400 ° C. However, if the cooling stop temperature is less than 200 ℃ not only saturation effect but also plate distortion due to over cooling can occur.

The steel sheet produced by the composition and the manufacturing method has a tensile strength of 930 MPa or more and an impact toughness of 200 J or more.

In the method described above, a fine structure of fine grained bainitic ferrite and acicular ferrite, which is effective for high strength and low temperature toughness, is formed, and thus it is possible to manufacture steel having excellent strength and unstable ductility resistance even in extreme regions. Tensile strength is 1.5 times higher than the most widely used tensile strength of 570 ~ 620MPa steel, but it can contribute to economic pipeline construction with excellent resistance to unstable ductile fracture.

The steel sheet produced by the composition and the manufacturing method is more than 85% DWTT ductile wavefront.

In the case of the pipeline constructed in the extreme area, in case of winter, the temperature of the pipeline drops to -40 ℃, and when the destruction occurs due to an accident, brittle fracture propagates rapidly and a large accident occurs in the pipeline. According to the manufacturing method of the present invention, when the mixed structure of the bainitic ferrite and the aciculal ferrite of the fine grains is formed, it is possible to secure more than 85% of the DWTT ductile fracture rate, which is a measure of suppressing the propagation of brittle fracture, thereby ensuring the stability of the pipeline. This is secured.

Hereinafter, the present invention will be described in detail.

EXAMPLE

The slab having the composition shown in Table 1 was heat-rolled-cooled under the production conditions shown in Table 2 to prepare a steel sheet having a thickness of 16 mm.

division C Si Mn Mo Cr Ni Ti Nb V Al Cu Ca * B * N * P * S * Inventive Steel A 0.050 0.26 1.91 0.30 0.30 1.00 0.014 0.050 0.040 0.020 0.3 15 12 41 70 12 Inventive Steel B 0.050 0.25 1.89 0.30 0.30 1.51 0.015 0.050 0.041 0.020 0.3 11 13 42 75 13 Invention Steel C 0.050 0.25 1.90 0.32 0.30 2.00 0.014 0.049 0.040 0.020 0.3 12 11 40 65 12 Comparative Steel D 0.050 0.25 1.91 0.30 0.30 0.50 0.014 0.050 0.040 0.019 0.3 13 12 41 68 15 Comparative Steel E 0.051 0.15 1.89 0.25 0.48 0.48 0.014 0.042 0.043 0.023 0.2 13 12 38 65 11 Comparative Steel F 0.050 0.15 2.31 0.25 0.50 1.01 0.015 0.040 0.040 0.020 0.2 10 20 30 80 15 Comparative Steel G 0.050 0.15 1.89 0.25 0.50 1.02 0.015 0.040 0.040 0.020 0.2 10 30 42 80 15

However, the content unit of the * element in the table is ppm, the content unit of the remaining elements is weight%.

As shown in Table 1, in the case of invention steel A to invention steel C, all of the conditions of the present invention are satisfied. Comparative steel D and comparative steel E correspond to the case where the Ni content is too low, and comparative steel F corresponds to the case where the Mn is excessively high. The comparative steel G also corresponds to the case where B is excessively high.

division Slab
Heating temperature
(℃) Recrystallization backpressure (%) Unresolved station
Rolling reduction
(%) Cooling rate
(℃ / sec) Cooling
Stop temperature
(℃) Rolling reduction
(%) Average rolling reduction
(%) Retention time
(second) Inventive Steel A Inventive Example 1 1110 6.3 12 <16 80 25 326 Inventive Steel A Inventive Example 2 1100 8.5 13 <14 90 30 230 Inventive Steel B Inventive Example 3 1125 6.0 12 <15 80 26 320 Inventive Steel B Inventive Example 4 1125 6.0 13 <15 90 28 272 Invention Steel C Inventive Example 5 1100 6.0 12 <15 80 25 316 Invention Steel C Inventive Example 6 1100 6.0 13 <15 90 29 250 Inventive Steel A Comparative Example 1 1100 6.3 14 <16 60 28 276 Inventive Steel A Comparative Example 2 1100 5.0 9 <16 80 29 250 Inventive Steel B Comparative Example 3 1125 6.3 14 <16 60 26 310 Invention Steel C Comparative Example 4 1100 6.3 14 <15 60 28 280 Comparative Steel D Comparative Example 5 1110 6.3 12 <16 80 26 300 Comparative Steel D Comparative Example 6 1110 6.0 13 <15 90 29 250 Comparative Steel E Comparative Example 7 1123 6.3 12 <16 80 25 343 Comparative Steel F Comparative Example 8 1118 6.3 12 <16 80 25 345 Comparative Steel G Comparative Example 9 1114 6.3 12 <16 30 28 289

Using a slab having the composition of Table 1, a portion of the steel sheet manufactured under the manufacturing conditions of Table 2 was taken, and the tensile strength, impact toughness, and DWTT ductile fracture rate were measured by performing a tensile test, an impact test, and a DWTT test. It is shown in Table 3 below.

Inventive Examples 1 to 6 of Table 2 satisfy all of the alloy composition and the manufacturing conditions of the present invention, and Comparative Examples 1 to 4 have the alloy composition of the inventive steel of Table 1, which is a composition satisfying the alloy composition of the present invention. It does not satisfy the manufacturing conditions. Comparative Examples 5 to 9 are cases where the production conditions of the present invention are applied to slabs having alloy compositions of Comparative Steels D to G in Table 1.

division Effective grain size
(Μm, based on the grain boundary 15 degrees) BF fraction
(%) The tensile strength
(MPa) Impact toughness
(J) DWTT
Ductility Rate (%) Inventive Steel A Inventive Example 1 5.7 53 972 225 95 Inventive Steel A Inventive Example 2 4.4 51.5 1050 241 100 Inventive Steel B Inventive Example 3 5.9 55.2 1031 230 92 Inventive Steel B Inventive Example 4 4.2 52.1 1039 226 95 Invention Steel C Inventive Example 5 6.0 57.2 1110 231 96 Invention Steel C Inventive Example 6 4.1 54.2 1055 220 97 Inventive Steel A Comparative Example 1 8.2 56.1 945 90 34 Inventive Steel A Comparative Example 2 9.3 54.7 955 67 46 Inventive Steel B Comparative Example 3 8.3 56.8 960 121 63 Invention Steel C Comparative Example 4 8.5 58.0 975 150 42 Comparative Steel D Comparative Example 5 5.4 43 940 35 34 Comparative Steel D Comparative Example 6 4.4 41 946 25 45 Comparative Steel E Comparative Example 7 5.5 64 955 34 25 Comparative Steel F Comparative Example 8 5.3 62.3 1060 121 50 Comparative Steel G Comparative Example 9 5.4 65.1 1070 68 46

However, the impact toughness is a value measured at -60 ° C and the DWTT ductility rate at -40 ° C, and BF means bainitic ferrite.

As shown in Table 3, in the case of the invention example having the composition and manufacturing conditions of the present invention, the tensile strength of 930 MPa or more, impact toughness of 200J or more at -60 ° C, or -40 ° C DWTT ductile wavefront ratio of 85% or more is a good value. Indicates.

Comparative Examples 1 to 4 have a composition that satisfies the conditions of the present invention, but do not satisfy the manufacturing conditions. Comparative Examples 1, 3, and 4 show a case where the unrecrystallized rolling reduction is low at 60%. The effective grain size is large due to the lack of unrecrystallized rolling reduction, and the impact is -60 ° C due to the high DBTT (ductility-brittle transition temperature). The toughness and -40 ° C DWTT characteristics were low. Comparative Examples 5, 6, and 7 have less Ni content than the present invention, and the effective grain size is finer than that of the present invention, but the bainitic ferrite content is low, resulting in low impact toughness and DWTT characteristics. In Comparative Examples 8 and 9, the Inventive Example and the Ni content were similar, but the Mn or B content was too high compared to the present invention, respectively, and the tensile strength was high, but the toughness was low.

In order to identify the cause of the excellent cryogenic toughness of the invention example compared to the comparative example, the processing-improved continuous cooling transformation properties of Inventive Example 1 and Comparative Example 5 were investigated and shown in FIGS. 1 and 2. As shown in Figures 1 and 2, the transformation of bainitic ferrite is encouraged as the Ni content is increased, so that even when the Ni content is manufactured under the same manufacturing conditions, the fraction of bainitic ferrite increases as the Ni content increases. Can be. FIG. 3 shows the results of scanning electron microscopy on the impact wave surface of Inventive Example 1, wherein cracks propagate zigzag along grain boundaries and lath boundaries of bainitic ferrite tissue. From this, it can be seen that the reason why the steel of the present invention retains high strength and high toughness at cryogenic temperatures is that a large amount of fine bainitic ferrite is formed at the time of rolling.

1 is a machining provision continuous cooling transformation diagram of Inventive Example 1. FIG.

Figure 2 is a machining provision continuous cooling transformation diagram of Comparative Example 5.

3 is a scanning electron micrograph of a shock wave surface according to Inventive Example 1. FIG.

Claims (3)

By weight%, C: 0.03-0.10%, Si: 0.05-0.6%, Mn: 1.6-2.1%, Cu: 0.5% or less (not containing 0%), Ni: 0.8-2.0%, Nb: 0.02- 0.06%, V: 0.1% or less (without 0%), Mo: 0.2-0.5%, Cr: 1.0% or less (without 0%), Ti: 0.005-0.03%, Al: 0.01-0.05 %, B: 0.0005 to 0.0020%, N: 0.002 to 0.006%, Ca: 0.0005 to 0.006%, P: 0.02% or less (without 0%), S: 0.005% or less (without 0%) , The remainder contains Fe and other unavoidable impurities, the microstructure consists of 51.5-57.2% bainitic ferrite and the remaining aciculous ferrite, cryogenic toughness characterized by a tensile strength of 930 MPa or more and an impact toughness of 200 J or more This super high strength steel sheet for line pipes. The steel sheet for ultra-high strength line pipe having excellent cryogenic toughness according to claim 1, wherein the DWTT ductile fracture rate of the steel sheet is 85% or more. delete

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