KR20080040233A - Steel plate for linepipe having ultra-high strength and excellent low temperature toughness and manufacturing method of the same - Google Patents

Abstract

A steel plate having high tensile strength and excellent low temperature toughness manufactured although Mn is not used in a large content is provided, and a manufacturing method of the steel plate is provided. A manufacturing method of a steel plate for a linepipe having ultra-high strength and excellent low temperature toughness comprises the steps of: heating a steel slab comprising, by weight percent, 0.03 to 0.10% of C, 0 to 0.6% of Si, 1.6 to 2.1% of Mn, 0 to 1.0% of Cu, 0 to 1.0% of Ni, 0.02 to 0.06% of Nb, 0.1% or less of V, 0.1 to 0.5% of Mo, 1.0% or less of Cr, 0.005 to 0.03% of Ti, 0.01 to 0.06% of Al, 0.0005 to 0.0025% of B, 0.001 to 0.006% of N, 0 to 0.006% of Ca, 0.02% or less of P, 0.005% or less of S, and the balance of Fe and other unavoidable impurities to a temperature of 1050 to 1150 deg.C; subjecting the heated slab to a single rolling operation or a multistage rolling operation of two or more rolling operations to a reduction ratio of 20 to 80% in a temperature range of an austenite recrystallization temperature or higher; subjecting the rolled slab to a single rolling operation or a multistage rolling operation of two or more rolling operations to a reduction ratio of 40 to 80% in a temperature range of the Ar3 temperature to the austenite recrystallization temperature to manufacture a steel plate; cooling the rolled steel plate at a cooling rate of 20 to 50 deg.C/sec; and stopping cooling of the steel plate at a temperature of 200 to 400 deg.C.

Description

Steel plate for ultra high strength line pipe with excellent low temperature toughness and manufacturing method thereof

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 steels composed of major microstructures of Lower Bainite and Lath Martensite and cooling conditions of steels consisting of major microstructures of Bainitic Ferrite and Acicular Ferrite; FIG.

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 a steel sheet for ultra high strength line pipe having excellent low temperature toughness and a method for manufacturing the same, and more particularly, a smaller amount of alloying element compared to the prior art in which a large amount of alloying element is added to secure strength and toughness. The present invention relates to a steel sheet having a strength of 930 MPa or more and excellent in toughness 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 the large diameter of the line pipe using this grade steel sheet inevitably requires an increase in the steel sheet thickness, which is not economical.

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 same structure such as martensite is formed inside the steel sheet, there is a problem that the toughness of the steel sheet is extremely poor due to insufficient strength of the steel sheet or residual stress inside the steel sheet. .

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 thermal history at the same time as mechanical processing for the steel sheet, but is used in a very large number of forms, mainly rolling under strict conditions 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 contain a large amount of Mo, in particular at least 0.2% by weight of Mo, in order to increase the strength of the steel, thereby reducing the steel internal structure to lower bainite and lath martensite. (Lath Martensite) is mainly described a technology, in particular, as a similar technology, referring to the Republic of Korea Patent No. 2000-00533890 as a main invention as well as a steel containing 0.15% by weight or more, as well as 0.14 weight It can be seen that the comparative example showing the Mo content of% shows a lower tensile strength than 930 MPa.

However, since Mo is an expensive element in itself, when Mo is used in the above-mentioned 0.15% by weight or 0.2% by weight or more, it causes a rise in steel manufacturing cost. In addition, the lower bainite structure shown in FIG. 3 has a very narrow temperature range in which phase transformation occurs, as shown in FIG. 2, so that the cooling conditions for producing steel are strict, and a very high cooling rate is required. As the capacity condition is very strict, as well as a high cooling rate, problems such as plate deformation may occur, and thus there is a problem that manufacturing conditions are cumbersome and difficult, such as additional processing for shape control after steel sheet production.

An object of the present invention is to provide a steel sheet having a high tensile strength and excellent low temperature toughness and a method of manufacturing the same without using a large amount of Mo and the like.

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 ~ 1.0 wt%, Nb: 0.02 to 0.06 wt%, V: 0.1 wt% or less, Mo: 0.1 to 0.5 wt%, Cr: 1.0 wt% or less, Ti: 0.005 to 0.03 wt%, Al: 0.01 to 0.06 wt%, B: 0.0005 to 0.0025% 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, residual Fe and other unavoidable impurities By adding the bainitic ferrite and the aciculus ferrite is characterized in that it comprises at least 75% by area fraction.

Steel sheet of another aspect of the present invention for achieving the above object is 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: 0.1 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 0.0025 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 It includes, including the bainitic ferrite and the acicular ferrite as the internal structure to include more than 75% based on the area fraction, characterized in that the tensile strength of 930MPa or more, -40 ℃ impact toughness of 230 Joules or more.

At this time, Mo is more preferably added 0.015% by weight or less.

In addition, the content of granular bainite is preferably 5% or less based on the area fraction.

And it is especially preferable that initial austenite grain size is 15 micrometers or less.

The production method of the present invention for achieving the above object is 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: 0.1 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 0.0025% 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 inevitable impurities Heating the slab 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 stopping the cooling of the steel sheet at a temperature of 200 to 400 ° C.

At this time, Mo is more preferably added at 0.015% by weight or less.

Moreover, it is effective to air-cool or to cool a steel plate after cooling stop of the said 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, as a result of reducing the content of Mo added to increase the strength of the steel sheet, even if the internal structure of the lower bainite or lath as conventional Even if it is not martensite, it is possible to exhibit sufficient strength and to secure low temperature toughness by finely adjusting the grain size during rolling and by adjusting the tissue separately from the superhard tissue such as lower bainite or lath martensite. It was confirmed that the present invention was reached.

That is, the present invention reduces the content of Mo and adjusts the amount of addition of other elements, while controlling by the bainitic ferrite and acicular ferrite (in other words acicular ferrite) having fine grains. A steel sheet having a strength equal to or higher than that of a steel sheet having a hard structure and having a low temperature toughness far superior to a conventional lower bainite or lath martensite by making the grain size of the structure fine; It is characterized by providing a special method.

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, this reinforcing effect cannot be obtained, and when it is added in excess of 0.1% by weight, the toughness of the base metal and the heat affected zone including the low temperature crack after welding is degraded.

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 development weldability and the toughness of the weld heat affected zone. Al or Ti plays a role of deoxidation, so it is not necessary to add Si for deoxidation.

Mn: 1.6 to 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.

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.

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%.

Mo: 0.1 ~ 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, the addition of excessive Mo lowers the toughness of the weld heat affected zone during the field welding, so it should be maintained at 0.5% by weight or less, and most preferably at 0.15% by weight or less.

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 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 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.

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 ~ 0.0025 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 0.0025% 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.

(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 the bainitic ferrite having the shape shown in FIG. 4 and the acyculous ferrite 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 addition, the steel sheet of the present invention needs to have a very fine internal structure. The finer the microstructure, the more it is possible to prevent the propagation of cracks to prevent brittle fracture. According to the inventors of the present invention, the preferred grain size is 15 µm or less based on the initial austenite grain size.

As described above, the steel sheet having a component system and satisfying the internal structure conditions is a steel sheet that satisfies all properties desired in the present invention as a tensile strength of 930 MPa or more and -40 ° C. impact toughness of 230 Joules or more.

(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, the slab is heated, and then 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 multi-stages at a temperature higher than the temperature (Ar ₃ ), the cooling is performed at a rate of 20 to 35 ° 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 the above 1050 ℃, Nb or V is not reusable in the steel, it is difficult to achieve high strength of the steel sheet, and partial recrystallization occurs, so that austenite grains are not uniformly formed, making it difficult to toughen it. 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, one roll or two or more multi-stage rollings are performed in the austenite recrystallization region at a total reduction ratio of 20 to 80% with respect to the initial slab thickness. 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. Fine austenite grains formed by the above-described process serves to improve the low temperature toughness 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, rolling is performed at a total reduction ratio of 40 to 80% with respect to the slab thickness that is rolled in the recrystallization temperature range. 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 a 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, when cooling at a high cooling rate of 50 ° C./sec or more, a hard phase such as martensite is formed due to the excessive amount of cooling water or a warpage of the steel sheet occurs, 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 to prepare a steel plate having a thickness of 16 mm. The manufacturing conditions of each steel plate were set the same. In other words, regardless of the type of slab, the heating temperature for the slab was 1120 ° C, followed by 9-9 steps of multi-step rolling at 73% of the total pressure reduction rate at 1050-1100 ° C (austenite recrystallization temperature), and then 700 9 to 11 stages of multi-stage rolling were performed at a total pressure reduction ratio of 76% at ˜950 ° C. (austenite recrystallization temperature). Immediately after rolling, cooling was performed at a cooling rate of 25 to 35 ° C./sec to terminate the cooling at 250 to 350 ° C., and then left to stand in air to allow air cooling.

division C Si Mn Mo Cr Ni Ti Nb V Al Cu Ca * B * N * P * S * Inventive Steel 1 0.049 0.15 1.89 0.14 0.47 0.47 0.014 0.037 0.037 0.025 0.2 14 16 43 54 12 Inventive Steel 2 0.051 0.15 1.89 0.13 0.50 0.50 0.015 0.040 0.040 0.023 0.2 10 10 39 70 10 Invention Steel 3 0.052 0.15 1.89 0.15 0.49 0.50 0.015 0.038 0.039 0.022 0.2 11 20 36 70 8 Inventive Steel 4 0.050 0.15 1.90 0.30 0.30 0.49 0.015 0.040 0.043 0.022 0.2 14 11 45 58 9 Comparative Steel 1 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 Comparative Steel 2 0.122 0.15 1.88 0.16 0.45 0.51 0.016 0.043 0.038 0.021 0.2 10 15 38 56 9 Comparative Steel 3 0.052 0.16 2.32 0.18 0.46 0.48 0.015 0.041 0.042 0.020 0.2 10 14 37 52 11 Comparative Steel 4 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 Comparative Steel 5 0.051 0.15 1.89 0.15 0.50 0.50 0.015 0.041 0.039 0.020 0.2 12 31 35 70 7 Comparative Steel 6 0.052 0.16 1.91 0.25 0.53 0.51 0.017 0.043 0.037 0.021 0.2 12 43 37 60 8

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 steels 1 to 4 are cases in which all of the conditions of the present invention are satisfied. However, comparative steel 1 corresponds to a case where C is too low and comparative steel 2 corresponds to a case where C is excessively high. In addition, comparative steel 3 is a case where Mn is excessively high, and comparative steel 4 is a case where Ti is excessively high. Comparative steels 5 and 6 correspond to the case where B is excessively high.

Table 2 shows the results of tensile test, impact test, and brittle fracture transition temperature by taking a portion of the steel sheet manufactured under the above conditions using the slab having the composition of Table 1.

division BF + AF fraction in internal tissues (%) Initial Austenitic Grain Size (㎛) Tensile Strength (MPa) vE-40 (Joule) vTrs (℃) Inventive Steel 1 83 12 984 276 -113 Inventive Steel 2 88 13 993 254 -94 Invention Steel 3 78 11 976 258 -93 Inventive Steel 4 92 13 965 245 -79 Comparative Steel 1 65 12 467 285 -98 Comparative Steel 2 86 13 1053 102 -48 Comparative Steel 3 88 12 1012 175 -64 Comparative Steel 4 84 11 983 187 -58 Comparative Steel 5 62 13 1034 192 -52 Comparative Steel 6 57 13 1038 136 -43

Here, vE-40 means impact toughness at -40 ° C, vTrs means brittle fracture transition temperature, BF means bainitic ferrite, and AF means acicular ferrite.

As can be seen from the results of Table 2, in the case of the invention steel having the composition of the present invention, the tensile strength is 930 MPa or more and impact toughness of 230 Joules or more at -40 ° C, and the brittle fracture transition temperature also shows a good value of -70 ° C or less. On the other hand, in the case of Comparative steel 1 having a low C content, the impact toughness was good, but the strength was too low as half the level of the inventive steel of the present invention. The strength was over 1000MPa, which showed ultra high strength, -40 ℃ impact energy was 102Joule, and the brittle fracture transition temperature was -48 ℃, indicating the problems of the conventional steel incompatible with strength and toughness. In addition, it can be confirmed that the comparative steel 3, in which Mn is excessively added, exhibits similar behavior to that of the comparative steel 2. Comparative steel 4 is a case where the amount of Ti is excessive, it can be seen that the impact energy and brittle fracture transition temperature is insufficient. In addition, it can be seen that the comparative steel 5 and the comparative steel 6, in which the B content is excessively high, are excellent in strength but not satisfactory in impact energy and brittle fracture transition temperature.

Therefore, the influence of the composition of the steel plate made into object by this invention was confirmed.

Among the steels satisfying the conditions of the present invention, a slab having a composition of Inventive Steel 1 was selected and rolled under the conditions of Table 3 below.

division Slab heating temperature (℃) Recrystallization rate reduction rate (%) Undetermined rolling reduction rate (%) Cooling rate (℃ / sec) Cooling stop temperature (℃) Inventive Example 1 1110 73 76 22 225 Inventive Example 2 1122 74 75 25 350 Inventive Example 3 1134 70 78 29 372 Inventive Example 4 1145 77 72 33 296 Comparative Example 1 1112 75 74 24 425 Comparative Example 2 1165 75 74 32 250 Comparative Example 3 1185 74 75 27 472 Comparative Example 4 1114 73 76 15 289 Comparative Example 5 1120 70 78 17 466 Comparative Example 6 1130 90 35 25 244

As can be seen from Table 3, Inventive Example 1 to Inventive Example 4 satisfy all of the conditions of the present invention, and Comparative Example 1 is a case where the cooling rate is excessively high. Comparative Example 2 and Comparative Example 3 is a case where the slab reheating temperature is excessively high, especially Comparative Example 3 shows a case where the slab heating temperature is excessively high and the cooling end temperature is excessively high. Comparative Example 4 and Comparative Example 5 is a case where the cooling rate is too low, especially Comparative Example 5 is a case where not only the cooling rate is too low but also the cooling end temperature is too high. Comparative Example 6 shows a case where the austenite unrecrystallized zone rolling reduction is too low.

The results of preparing the steel sheets in the respective manners described in Table 3 are shown in Table 4 below.

division BF + AF fraction in internal tissues Initial Austenitic Grain Size Tensile Strength (MPa) vE-40 (Joule) vTrs (℃) Inventive Example 1 88 8 984 276 -113 Inventive Example 2 84 12 993 254 -94 Inventive Example 3 88 14 976 258 -93 Inventive Example 4 94 13 965 245 -79 Comparative Example 1 65 14 865 268 -98 Comparative Example 2 86 25 934 197 -64 Comparative Example 3 88 30 823 189 -50 Comparative Example 4 64 15 834 187 -48 Comparative Example 5 61 12 852 218 -52 Comparative Example 6 57 15 879 143 -42

Here, vE-40 means impact toughness at -40 ° C, vTrs means brittle fracture transition temperature, BF means bainitic ferrite, and AF means acicular ferrite.

As can be seen in Table 4, in the case of Inventive Examples 1 to 4, which satisfy all the conditions of the present invention, it was confirmed that all of them had excellent properties of tensile strength of 930 MPa or more and impact energy (-40 ° C.) of 230 J or more. . However, in Comparative Example 1, the cooling end temperature was too high, and thus the fine low-temperature phases could not be formed properly, resulting in low tensile strength. In Comparative Example 2, the slab reheating temperature was excessively high, resulting in low temperature toughness due to coarse austenite grains. In Comparative Example 3, the slab reheating temperature was excessively high and the cooling end temperature was too high. Thus, the low-temperature toughness was low for the same reason as in Comparative Example 2, and the tensile strength was low for the same reason as in Comparative Example 1. Comparative Example 4 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 low temperature toughness. In Comparative Example 5, the cooling rate was too low and the cooling end temperature was excessively high. Thus, both the tensile strength and the low temperature toughness were low. Comparative Example 6 shows that the austenite unrecrystallized zone reduction ratio is too low, and the austenite grains are not drawn properly, and dislocations do not accumulate inside the grains, and thus low temperature toughness is not formed properly. It was.

Therefore, the effect of the manufacturing method which concerns on this invention was confirmed.

As described above, according to the present invention, it is possible to provide a steel sheet for line pipe having excellent low temperature toughness as well as ensuring high strength without adding a large amount of Mo or the like.

Claims (9)

By weight% C: 0.03 ~ 0.10 wt%, Si: 0 ~ 0.6 wt%, Mn: 1.6 ~ 2.1 wt%, Cu: 0 ~ 1.0 wt%, Ni: 0 ~ 1.0 wt%, Nb: 0.02 ~ 0.06 wt% , V: 0.1 wt% or less, Mo: 0.1-0.5 wt%, Cr: 1.0 wt% or less, Ti: 0.005-0.03 wt%, Al: 0.01-0.06 wt%, B: 0.0005-0.0025 wt%, N: 0.001 ~ 0.006% by weight, Ca: 0 ~ 0.006% by weight, P: 0.02% by weight or less, S: 0.005% by weight or less, including residual Fe and other unavoidable impurities, and combining bainitic ferrite and acicular ferrite as an internal structure Steel sheet for ultra high strength line pipe having excellent low temperature toughness, characterized in that it comprises 75% or more based on the area fraction. By weight% C: 0.03 ~ 0.10 wt%, Si: 0 ~ 0.6 wt%, Mn: 1.6 ~ 2.1 wt%, Cu: 0 ~ 1.0 wt%, Ni: 0 ~ 1.0 wt%, Nb: 0.02 ~ 0.06 wt% , V: 0.1 wt% or less, Mo: 0.1-0.5 wt%, Cr: 1.0 wt% or less, Ti: 0.005-0.03 wt%, Al: 0.01-0.06 wt%, B: 0.0005-0.0025 wt%, N: 0.001 ~ 0.006% by weight, Ca: 0 ~ 0.006% by weight, P: 0.02% by weight or less, S: 0.005% by weight or less, including residual Fe and other unavoidable impurities, and combining bainitic ferrite and acicular ferrite as an internal structure The steel sheet for ultra-high strength line pipe having excellent low-temperature toughness, including 75% or more by area fraction, having a tensile strength of 930MPa or more and -40 ° C impact toughness of 230 Joules or more. The steel sheet for ultra-high strength line pipe having excellent low temperature toughness according to claim 1 or 2, wherein Mo is added at 0.015% by weight or less. The steel sheet for ultra-high strength line pipe having excellent low temperature toughness according to claim 1 or 2, wherein the content of granular bainite is 5% or less based on the area fraction. The steel sheet for ultra-high strength line pipe according to claim 1 or 2, wherein the old austenite grain size is 15 µm or less. The steel sheet for ultra-high strength line pipe having excellent low temperature toughness according to claim 1 or 2, wherein the steel sheet has a tensile strength of 930 MPa or more and an impact toughness of 230 Joules or more at -40 ° C. By weight% C: 0.03 ~ 0.10 wt%, Si: 0 ~ 0.6 wt%, Mn: 1.6 ~ 2.1 wt%, Cu: 0 ~ 1.0 wt%, Ni: 0 ~ 1.0 wt%, Nb: 0.02 ~ 0.06 wt% , V: 0.1 wt% or less, Mo: 0.1-0.5 wt%, Cr: 1.0 wt% or less, Ti: 0.005-0.03 wt%, Al: 0.01-0.06 wt%, B: 0.0005-0.0025 wt%, N: 0.001 ~ 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, steel slabs containing residual 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 Stopping the cooling of the steel sheet at a temperature of 200 ~ 400 ℃; Method for producing a steel sheet for ultra-high strength line pipe having excellent low temperature toughness, characterized in that it comprises a. 8. The method of claim 7, wherein the Mo is added at 0.015% by weight or less. The method for manufacturing an ultra-high strength line pipe sheet having excellent low-temperature toughness according to claim 7 or 8, wherein the steel sheet is cooled by air or by cooling after the cooling stop of the steel sheet.

Images