KR100843844B1 - Steel plate for ultra high strength line pipe with excellent crack growth resistance and manufacturing method - Google Patents

Abstract

The present invention relates to an ultra-high strength line pipe sheet having excellent crack growth resistance and a method for manufacturing the same, and more particularly, to precisely control the rolling method in order to secure strength and crack growth resistance, and having a strength of 930 MPa or more and having crack growth resistance. It relates to an excellent steel sheet and a method of manufacturing the same.

The manufacturing method of the steel sheet for ultra-high strength line pipe having excellent crack growth resistance is C: 0.03 to 0.10 wt%, Si: 0 to 0.6 wt%, Mn: 1.6 to 2.1 wt%, Cu: 0 to 1.0 wt% , Ni: 0 to 1.0% by weight, Nb: 0.02 to 0.06% by weight, V: 0.1% by weight or less, Mo: 0.2 to 0.5% by weight, Cr: 1.0% by weight or less, Ti: 0.005 to 0.03% by weight, Al: 0.01 ~ 0.06 wt%, B: 0.0005 to 0.0020 wt%, N: 0.001 to 0.006 wt%, Ca: 0 to 0.006 wt%, P: 0.02 wt% or less, S: 0.005 wt% or less, balance Fe and other unavoidable impurities Heating the comprising steel slab to a temperature of 1050-1150 degrees Celsius; Rolling the heated slab once or multi-stage two or more times at a reduction ratio of 20 to 80% in a temperature range of at least austenite recrystallization temperature; Manufacturing the rolled slab by rolling one or two or more times in multiple stages at a reduction ratio of 40 to 80% in a temperature range of less than or equal to austenite recrystallization temperature, Ar3; Cooling the rolled steel sheet at a cooling rate of 20 to 50 ° C./sec; And stopping the cooling of the steel sheet at a temperature of 200 ° C. to 400 ° C., wherein a rolling reduction rate of each rolling step in a temperature section of the austenite recrystallization temperature or more is 5% or more, and the temperature of the austenite recrystallization temperature or more. The average reduction ratio (the value obtained by dividing the total reduction ratio for each step by the number of reduction steps) in the section is 10% or more.

Bainitic ferrite, high strength, high toughness, DWTT, linepipe

Description

Steel plate for ultra-high strength line pipe with excellent crack growth resistance and its manufacturing method {STEEL PLATE FOR LINEPIPE HAVING ULTRA-HIGH STRENGTH AND EXCELLENT CRACK PROPAGATION RESISTANCE AND MANUFACTURING METHOD OF THE SAME}

1 is a schematic graph comparing a manufacturing method of tempering to secure properties of steel produced through rolling and cooling and a manufacturing method of securing physical properties without tempering;

FIG. 2 is a TTT diagram for comparing cooling conditions of steel composed of major microstructures of lower bainite and lath martensite with cooling conditions of steel composed of major microstructures of bainitic ferrite and aciculous ferrite;

3 is a transmission electron microscope observation photograph of the lower bainite;

4 is a transmission electron microscope observation photograph of bainitic ferrite;

5 is a transmission electron microscope observation picture of the acicular ferrite; And

6 is a transmission electron microscope observation photograph of granular bainite.

The present invention relates to an ultra-high strength line pipe sheet having excellent crack growth resistance and a method for manufacturing the same, and more particularly, to precisely control the rolling method in order to secure strength and crack growth resistance, and having a strength of 930 MPa or more and having crack growth resistance. It relates to an excellent steel sheet 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, simply referred to as 'crude oil') that can be transported per unit time. For this purpose, the diameter of the line pipe is necessarily increased to a large diameter. I need to.

When the diameter of the line pipe is increased, the amount of crude oil or the like flowing therein may be increased, and thus the pressure acting on the line pipe by crude oil is also increased. In this case, the material for the line pipe needs to be developed with higher strength. Until now, it is a reality that steel sheets of grade X70 are mainly used in view of the strength standards of the line pipe. The X70 grade steel sheet has a strength of about 70 ksi, that is, about 480 MPa, and it is inevitably required to increase the thickness of the steel sheet to large diameter of the line pipe using such strength grade steel sheet.

Therefore, the demand for development of a steel sheet whose strength is dramatically improved as compared to the line pipe steel sheet commonly used until now, but the development of a steel sheet that completely meets these requirements has not been completed yet. It is a reality.

The reason for this is that not only the barrier to the technique itself for increasing the strength of the steel sheet but also other problems caused by increasing the strength of the steel sheet is difficult to apply.

In other words, in order to increase the strength of the steel sheet, an alloying element effective for increasing the strength should be added, and it is difficult to obtain a sufficiently high strength by the addition of such an alloying element. Since the low temperature toughness and the low temperature toughness of the base material are also very poor, it is necessary to improve the low temperature toughness at the same time when high strength steel sheet.

In addition, in order to improve the strength of the steel sheet, conventionally, the steel sheet is quenched to form a low-temperature structure, particularly a lower bainite or martensite, inside the steel sheet, thereby increasing the hardness of the steel sheet and improving the strength thereof. When the martensitic structure is formed inside the steel sheet, there is a problem that the toughness of the steel sheet is extremely poor due to the residual stress in the steel sheet.

In particular, the method for evaluating the toughness required for line pipe steel is divided into a Charpy impact test indicating crack initiation resistance and a Drop Weight Tear Test (DWTT) indicating crack growth resistance. The ultra-high strength steel of API-X100 grade or more developed previously has poor crack growth resistance, which is evaluated as DWTT, and it is designed to additionally install a crack arrestor to suppress crack growth during pipeline construction. This increases the cost for pipeline construction and offsets the economic benefits of high strength steel.

As described above, the strength and toughness of the steel sheet are two conventionally incompatible physical properties. As the strength of the steel sheet increases, the general recognition is that the toughness decreases.

Since then, efforts have been made to provide high strength steels by securing the strength and toughness of the steel sheet at the same time. As one of the methods, a method called TMCP (Thermo Mechanical Controlling Process) has been proposed. The TMCP method is a general term for the processing method for changing the physical properties of the steel sheet to the desired physical properties by applying a thermal history at the same time as the mechanical processing for the steel sheet, it is used in very many forms, but mainly rolling under a strict condition at a predetermined temperature It consists of a controlled rolling process and an accelerated cooling process that cools the steel plate at an appropriate range of cooling rates.

When using the TMCP method, there is an advantage in that the desired physical properties can be smoothly controlled to a certain extent in theory by forming fine grains in the steel sheet and controlling the structure in a desired form.

However, in order to produce a steel sheet having a desired strength through the accelerated cooling of the TMCP as described above, it is necessary to form a hard structure as in the prior art. Therefore, even in the case of steel sheet manufactured by the TMCP method, it is inevitable in the prior art that the toughness is in the tendency to decrease as the strength increases.

Therefore, in the field of high strength steel, efforts have been made to secure a means for securing low temperature toughness while continuously conducting research and development to increase the strength level of steel.

In order to solve this problem, various methods have been proposed. Tempering on the manufactured steel sheet is the most widely used method.

For example, U.S. Patents 5545269, 5755895, 5798004, 5900075, 6045630, 6183573, 6245290, 6532995 have been subjected to TMCP treatment to complete the rolling and cooling as shown in Figure 1 Ac ₁ transformation temperature (ferrite is austenite when heated) The manufacturing method includes a process of additionally performing tempering under the temperature of transformation into a knight.The energy consumption is large and tempering because the steel sheet needs to be reheated to perform tempering after cooling again after cooling the steel sheet. Since the process must be added separately, there is a problem that the cost rises.

Further, in order to increase the strength of the steel, various kinds of alloying elements are added to the steel, and Mo is one of the most effective elements. For example, U.S. Pat.Nos. 6224689, 6228183, 6248191 and 6264760 dramatically improve the hardenability by adding B in combination with Mo to reduce the lower internal bainite and lath martensite. Techniques that mainly include are described.

However, all of the above techniques deal only with crack initiation resistance from the viewpoint of low temperature toughness, and there is little known technique for improving crack growth resistance for ultra high strength steel of API-X100 steel or higher.

An object of the present invention is to provide a method for manufacturing a steel sheet having high tensile strength and excellent crack growth resistance by precisely controlling the rolling method.

Method for producing a steel sheet of the present invention for achieving the above object by weight% C: 0.03 ~ 0.10% by weight, Si: 0 ~ 0.6% by weight, Mn: 1.6 ~ 2.1% by weight, Cu: 0 ~ 1.0% by weight, Ni : 0 to 1.0% by weight, Nb: 0.02 to 0.06% by weight, V: 0.1% by weight or less, Mo: 0.2 to 0.5% by weight, Cr: 1.0% by weight or less, Ti: 0.005 to 0.03% by weight, Al: 0.01 to 0.06 % By weight, B: 0.0005 to 0.0020% by weight, N: 0.001 to 0.006% by weight, Ca: 0 to 0.006% by weight, P: 0.02% by weight or less, S: 0.005% by weight or less, balance Fe and other unavoidable impurities Heating the steel slab to a temperature of 1050-1150 degrees Celsius; Rolling the heated slab once or multi-stage two or more times at a reduction ratio of 20 to 80% in a temperature range of at least austenite recrystallization temperature; Manufacturing the rolled slab by rolling one or two or more times in multiple stages at a reduction ratio of 40 to 80% in a temperature range of less than or equal to austenite recrystallization temperature, Ar3; Cooling the rolled steel sheet at a cooling rate of 20 to 50 ° C./sec; And stopping the cooling of the steel sheet at a temperature of 200 ° C. to 400 ° C., wherein a rolling reduction rate of each rolling step in a temperature section of the austenite recrystallization temperature or more is 5% or more, and the temperature of the austenite recrystallization temperature or more. In the interval, the average reduction ratio (the value obtained by dividing the total reduction ratio for each step by the number of reduction stages) is 10% or more.

At this time, it is effective to have a manufacturing condition in which the holding time between each pressing step is 20 seconds or less.

And after cooling stop of the said steel plate, it is preferable to air-cool or to cool a steel plate.

Hereinafter, the present invention will be described in detail.

The inventors of the present invention have studied in depth the method for solving the problems of the prior art, so that even if the internal structure is not lower bainite or lath martensite as in the prior art, while showing sufficient strength and rolling at the same time It was confirmed that the early austenite grain size can be finely adjusted to ensure crack growth resistance at the time of the present invention.

That is, the present invention is controlled by the bainitic ferrite (Bainitic Ferrite) and the acicular ferrite (in other words acicular ferrite) having a fine grain has a strength equal to or higher than that of a steel sheet having a conventional hard tissue, It is characterized by providing a steel sheet which has a much better crack growth resistance than conventional lower bainite or lath phase martensite by making the grain size of the structure fine, and a special method of manufacturing the steel sheet.

Hereinafter, the conditions of the present invention for achieving the above object will be described in detail in the order of the composition, internal structure and manufacturing method of the steel sheet of the steel sheet.

(Composition of steel sheet)

In the present invention, the composition of the steel sheet as a target was selected as follows in order to have sufficient strength and welded part toughness.

C: 0.03 ~ 0.10 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 plays a role of improving the hardenability, which is the ability 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 ~ 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. Excessive addition of Si to 0.6% by weight greatly reduces the field weldability and the toughness of the weld heat affected zone. Since Al or Ti plays a role of deoxidation, it is not necessary to add Si for deoxidation.

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 ~ 1.0 wt%

Cu is an element to strengthen the base metal and the weld heat affected zone. However, excessively added Cu lowers the toughness and field weldability of the weld heat affected zone.

Ni: 0 ~ 1.0 wt%

Ni is an element that improves physical properties without losing on-site weldability and low temperature toughness in low carbon steels. 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, it suppresses the occurrence of surface cracking in Cu-added steel during continuous casting and hot rolling. However, Ni is an expensive element and excessive addition of Ni lowers the toughness of the weld heat affected zone. More preferably, 0.2 to 0.8 wt% is added.

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 be contained 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 0.02-0.08 weight%.

V: 0.1 wt% or less

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.

Cr: 1.0 wt% or less

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 the heating process for rolling, and also reacts with nitrogen remaining in Ti. The solid solution in the steel is combined with carbon to form TiC precipitates, 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 0.03% by weight or more is added, coarsening of Ti nitride and 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 ~ 0.06 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 0.06% by weight, Al oxide (Al ₂ O ₃ ) is formed to reduce the toughness of the base metal and the heat affected zone. Since deoxidation can be carried out through the addition of Ti and Si, Al is not necessarily added.

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. 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: 0.001 ~ 0.006 wt%

N suppresses austenite grain growth during slab heating, and TiN precipitate suppresses austenite grain growth of the weld heat affected zone. Excessive N additions, however, promote slab surface defects and reduce the hardenability effect of B and, in the presence of nitrogen, degrade the toughness of the matrix and weld heat affected zones.

Ca: 0 ~ 0.006 wt%

Ca is mainly used as an element to control the shape of MnS inclusions and to improve low temperature toughness. However, excessive Ca addition causes a large amount of CaO-CaS to form and combine to form coarse inclusions, thereby degrading the cleanliness of the steel and damaging the field weldability.

P: less than 0.02% by weight

P is combined with Mn to form a non-metallic inclusion to cause the problem of embrittlement of the steel, so it is necessary to actively reduce it.However, in order to reduce P to an extreme, the steelmaking process load is intensified and the above problem occurs significantly at 0.02 wt% or less. Since it is not, the upper limit is made into 0.02 weight%.

S: 0.005 wt% or less

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.

(Internal organization)

As the steel sheet having the above-described component system, it is necessary to further limit the type and shape of the internal structure as preferable conditions for forming a steel sheet having excellent strength and excellent toughness.

That is, the microstructure inside the steel sheet provided in the present invention should contain at least 75% of bainitic ferrite having the shape shown in FIG. 4 and acyclic currite having the shape shown in FIG. 5. Here, the ratio of the tissue means an area fraction.

Some granular bainite may be formed in addition to the internal tissue of the above type. Since the granular bainite is a cause of inhibiting low-temperature toughness, its content should be limited to 5% or less on an area fraction basis.

In particular, the steel sheet of the present invention needs to have a very fine internal structure. The finer the microstructure, the better the crack growth resistance. According to the inventors of the present invention, the preferred grain size should be in the range of 15 μm based on the initial austenite grain size, and more preferably 10 μm or less. .

As described above, the steel sheet having the component system and satisfying the internal structure conditions is a steel sheet that satisfies all properties desired in the present invention as the DWTT ductile wave rate at -20 ° C. is greater than or equal to 930 MPa.

(Manufacturing method)

The most preferred method elicited by the present inventors for producing the steel which satisfies the object of the present invention as described above is described below.

In the manufacturing method of the present invention, after the slab is heated, the heated slab is subjected to one or two or more steps of multi-stage rolling in an austenite recrystallization zone, and then the austenite is lower than the austenite recrystallization temperature and transformed into a ferrite. After finishing rolling in one or two or more stages at a temperature higher than the temperature (Ar ₃ ), the cooling is performed at a rate of 20 to 50 ° C./sec and the cooling is terminated at 200 to 400 ° C. 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.

Slab heating: 1050 ~ 1150 ℃

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 the heating temperature is required. 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 1150 ° C., the austenite grains are excessively coarse, thereby providing a cause of increasing the grain size of the steel sheet, and as a result, the toughness of the steel sheet is extremely deteriorated. Therefore, the appropriate heating temperature range is preferably 1050 ~ 1150 ℃.

Rolling condition

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, 20-80% 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-described process serves to improve the crack growth resistance of the final plate. After this, one rolling or two or more multistage rollings are performed in an 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, 40 to 80% of the slab thickness after rolling in the recrystallization temperature range is performed. The 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

Cooling Speed: 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 slow Polygonal Ferrite or granular bay of the type shown in FIG. Undesirable tissues such as nitrate (Granular Bainite) are formed with coarse grain size, which may greatly reduce the strength and toughness to a level that does not meet the target of the present invention. On the contrary, in the case of cooling at a high cooling rate of 50 ° C./sec or more, warpage of the steel sheet occurs due to the excessive amount of cooling water, resulting in poor shape of the steel sheet.

Cooling end: 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 is stopped, is more than 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, when the cooling stop temperature is less than 200 ℃ not only saturation effect but also plate distortion due to excessive cooling may occur.

(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. In the recrystallization temperature range, multistage rolling was performed under the conditions shown in Table 2.

division C Si Mn Mo Cr Ni Ti Nb V Al Cu Ca * B * N * P * S * Invention steel A 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 B 0.058 0.16 1.88 0.27 0.52 0.52 0.017 0.041 0.042 0.022 0.2 11 14 42 71 13 C 0.048 0.15 1.90 0.32 0.49 0.51 0.015 0.039 0.039 0.021 0.2 12 20 34 60 9 D 0.060 0.16 1.80 0.26 0.41 0.53 0.016 0.045 0.043 0.024 0.2 14 18 46 76 15 Comparative steel E 0.024 0.16 1.91 0.25 0.38 0.51 0.016 0.041 0.042 0.020 0.2 13 18 39 65 11 F 0.134 0.16 1.88 0.36 0.45 0.51 0.016 0.043 0.041 0.021 0.2 10 15 41 66 13 G 0.053 0.16 2.32 0.28 0.46 0.46 0.015 0.046 0.040 0.023 0.2 11 17 36 72 16 H 0.051 0.15 1.88 0.20 0.46 0.49 0.038 0.042 0.040 0.020 0.2 12 14 42 62 10 I 0.061 0.15 1.89 0.35 0.40 0.50 0.015 0.047 0.042 0.020 0.2 12 31 39 78 9

However, the content units of the * elements in the table is ppm, and the content units of the remaining elements are weight percent.

As can be seen from Table 1, the invention steel A to the invention steel D is a case of satisfying all the conditions of the present invention, and the comparative steel E to the invention steel I is a case out of the conditions of the present invention. Comparative steel E corresponds to a case where C is too low and comparative steel F corresponds to a case where C is excessively high. In addition, comparative steel G is the case where Mn is excessively high, and comparative steel H is the case where Ti is excessively high. Comparative steel I corresponds to the case where B is excessively high.

division Slab heating temperature (℃) Recrystallization backpressure (%) Undetermined rolling reduction rate (%) Cooling rate (℃ / sec) Cooling stop temperature (℃) Rolling reduction (%) Average rolling reduction rate (%) Rolling steps Retention time (seconds) Inventive Example A One 1110 6.3 12 11 <16 73 25 326 B One 1122 10.2 15 9 <14 75 34 383 C One 1134 8.5 12 11 <9 76 23 276 D One 1145 7.1 14 11 <12 72 33 297 Comparative example A 2 1168 7.5 15 9 <10 74 32 250 A 3 1112 2.1 12 11 <14 75 34 253 A 4 1120 6.8 9 13 <9 75 28 323 A 5 1142 6.9 14 11 <23 76 35 287 A 6 1123 9.4 17 9 <40 75 29 343 A 7 1118 8.6 16 16 <11 34 31 345 A 8 1114 6.8 15 9 <19 76 15 289 A 9 1125 6.3 13 11 <13 74 24 468 E One 1120 7.6 17 9 <14 75 27 326 F One 1120 7.6 17 9 <14 75 33 293 G One 1120 7.6 17 9 <14 75 26 267 H One 1120 7.6 17 9 <14 75 33 257 I One 1120 7.6 17 9 <14 75 35 345

However, the holding time in the table is the time to be air-cooled between the rolling step.

Table 3 shows the results of measuring tensile strength and DWTT ductility by performing a tensile test and a DWTT test by taking a portion of the steel sheet manufactured according to the manufacturing conditions of Table 2 using the slab having the composition of Table 1 above. Inventive Examples A1 to D1 of Table 2 satisfy all of the alloy composition and manufacturing conditions of the present invention, and Comparative Examples A2 to A9 have an alloy composition of Inventive Steel A of Table 1, which is a composition that satisfies the alloy composition of the present invention. This is the case where the manufacturing conditions of the invention are not satisfied. In addition, Comparative Examples E1 to I1 are cases where the production conditions of the present invention are applied to the slab having the alloy composition of Comparative Steels E to I in Table 1.

division Initial Austenitic Grain Size (㎛) BF + AF fraction (%) Tensile Strength (MPa) DWTT Ductility Rate (%) Invention steel A One 12.7 83 984 95 B One 9.5 88 993 100 C One 10.5 78 976 95 D One 13.2 92 965 88 Comparative steel A 2 23.5 87 956 34 A 3 19.3 77 897 46 A 4 18.8 78 945 63 A 5 25.5 79 975 42 A 6 28.3 76 879 34 A 7 13.3 57 879 45 A 8 12.5 64 836 87 A 9 14.5 62 856 89 E One 12.2 65 467 100 F One 13.7 87 1058 35 G One 12.5 88 1022 41 H One 11.3 84 983 49 I One 13.5 64 1025 40

However, the DWTT ductility is the result at -20 ° C, BF means bainitic ferrite and AF means acicular ferrite.

As can be seen from the results of Table 3, in the case of the invention examples having the composition and manufacturing conditions limited in the present invention, both the tensile strength of 930 MPa or more, DWTT ductile wavefront at -20 ℃ 85% or more showed good values have.

Comparative Examples A2 to A9 have a composition that satisfies the conditions of the present invention, but are conditions that do not satisfy the manufacturing conditions. A2 is a case where the slab heating temperature is excessively high. The grain size of the austenite when extracted from the heating furnace is coarse, and the size of the austenite grain is not fine even after rolling in the austenite recrystallization region. Too low. A3 has a very low minimum reduction ratio in the austenite recrystallization region, and A4 has a very low average reduction ratio in the austenitic recrystallization region. In both cases, DWTT ductile wavefront ratio is too high because fine austenite grains cannot be formed. Low values were shown. A5 and A6 are cases where the austenite grain growth occurs because the holding time between the rolling steps is too long, and the austenite grain size is coarse, and the DWTT ductile wavefront ratio is too low. A7 is a case where the total reduction ratio in the austenite unrecrystallized region is too low, and the austenite grains are not drawn properly, and the dislocations do not accumulate inside the grains, and thus the low temperature transformation phase is not properly formed so that the tensile strength and the DWTT soft wave The results were very low. A8 is a case where the cooling rate is too low to properly form the tissue intended in the present invention, polygonal ferrite or granular bainite, etc. are mixed, showing both low tensile strength and DWTT ductility. A9 is a case where the cooling end temperature is too high, the fine low-temperature phases are not formed properly, resulting in low tensile strength.

E1 to I1 have manufacturing conditions satisfying the conditions of the present invention, but the composition does not satisfy the conditions of the present invention. In the case of E1 with low C content, the DWTT ductility is good but the tensile strength is too low as 1/2 level of the inventive steel, and in the case of F1 with excessively high C content, the tensile strength is over 1000 MPa. As a result, it showed very high strength and the DWTT ductility was very low. In addition, it can be seen that G1 added with excessively Mn exhibits similar behavior to F1. H1 is a case where the amount of Ti added is excessive, it can be seen that the DWTT ductile fracture rate is insufficient. In addition, in the case of I1, which was excessively high in B content, the strength was excellent, but the DWTT ductility was not satisfactory.

Therefore, the effect of the alloy composition and manufacturing method according to the present invention was confirmed.

As described above, according to the present invention, it is possible to provide a line pipe steel sheet having excellent crack growth resistance while securing ultra high strength by controlling the rolling in the austenite recrystallization region to form fine initial austenite.

Claims (3)

C: 0.03 to 0.10% by weight, Si: 0 to 0.6% by weight, Mn: 1.6 to 2.1% by weight, Cu: 0 to 1.0% by weight (not including 0%), Ni: 0 to 1.0% by weight (Not including 0%), Nb: 0.02 ~ 0.06% by weight, V: 0.1% by weight or less (0%), Mo: 0.2 ~ 0.5% by weight, Cr: 1.0% by weight or less (0% Ti: 0.005 ~ 0.03% by weight, Al: 0.01 ~ 0.06% by weight, B: 0.0005 ~ 0.0020% by weight, N: 0.001 ~ 0.006% by weight, Ca: 0 ~ 0.006% by weight (not including 0%) Steel slab containing P: 0.02 wt% or less, S: 0.005 wt% or less, balance Fe and other unavoidable impurities, Heating to a temperature of 1050-1150 ° C .; Rolling the heated slab once or multi-stage two or more times at a reduction ratio of 20 to 80% in a temperature range of at least austenite recrystallization temperature; Manufacturing the rolled slab by rolling one or two or more times in multiple stages at a reduction ratio of 40 to 80% in a temperature range of less than or equal to austenite recrystallization temperature, Ar3; Cooling the rolled steel sheet at a cooling rate of 20 to 50 ° C./sec; And To stop the cooling of the steel sheet at a temperature of 200 ~ 400 ℃; includes, The rolling reduction rate of each rolling step in the temperature section above the austenite recrystallization temperature is 5% or more, and the average rolling reduction rate (the total of the reduction rates in each step divided by the number of reduction steps) in the temperature section above the austenite recrystallization temperature is A method for producing a steel sheet for ultra high strength line pipe, which has excellent crack growth resistance, characterized in that more than 10%. The method for manufacturing an ultra-high strength line pipe sheet having excellent crack growth resistance according to claim 1, further comprising a production condition in which the holding time between the reduction steps in the austenite recrystallization region is 20 seconds or less. The method for manufacturing an ultra-high strength line pipe sheet according to claim 1 or 2, wherein the steel sheet is cooled by air or by cooling after cooling stop of the steel sheet.

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