WO2011030768A1

WO2011030768A1 - Steel sheet for high-strength line pipe having excellent low-temperature toughness, and steel pipe for high-strength line pipe

Info

Publication number: WO2011030768A1
Application number: PCT/JP2010/065354
Authority: WO
Inventors: 卓也原; 泰志藤城; 好男寺田; 豪鈴木; 太郎村木
Original assignee: 新日本製鐵株式会社
Priority date: 2009-09-09
Filing date: 2010-09-01
Publication date: 2011-03-17
Also published as: JP5131715B2; CN102482744B; BR112012005189A2; CN102482744A; JPWO2011030768A1

Abstract

Disclosed are a steel sheet for a line pipe and a steel pipe for a line pipe, which have excellent low-temperature toughness and are most suitable for a steel pipe that is used for a line pipe for transportation of petroleum oil, natural gas or the like. The steel sheet for a line pipe and the steel pipe for a line pipe respectively contain, in mass%, 0.02-0.08% of C, 0.01-0.5% of Si, 1.2-1.8% of Mn, 0.001-0.10% of Nb, 0.0010-0.0050% of N, and 0.0001-0.0050% of Ca, while suppressing P to 0.01% or less, S to 0.0020% or less, Ti to 0.001-0.030%, Al to 0.030% or less and O to 0.0035% or less with the balance made up of Fe and unavoidable impurities. The steel sheet for a line pipe and the steel pipe for a line pipe satisfy the relation of S/Ca < 0.5, and have a maximum Mn segregation degree of 2.0 or less, an Nb segregation degree of of 4.0 or less and a Ti segregation degree of 4.0 or less.

Description

Steel plate for high-strength line pipe and steel pipe for high-strength line pipe with excellent low-temperature toughness

The present invention relates to a steel plate for a line pipe and a steel pipe for a line pipe, which are excellent in low temperature toughness and are suitable for applications such as oil and natural gas transportation line pipes.

In recent years, pipelines have become increasingly important as long-distance transportation methods for crude oil and natural gas. Currently, the American Petroleum Institute (API) standard X65 is the basic design for trunk line pipes for long-distance transportation, and the actual usage is overwhelmingly large. However, a higher-strength line pipe is required for (1) improving transportation efficiency by increasing pressure and (2) improving local construction efficiency by reducing the outer diameter and weight of the line pipe. Up to now, line pipes up to X80 (tensile strength of 620 MPa or more) have been put into practical use, but the need for higher-strength line pipes has become stronger. Currently, research on high-strength line pipe manufacturing methods includes X80 line pipe manufacturing technology (Non-Patent Documents 1 and 2), X100 (tensile strength of 760 MPa or more) line pipe manufacturing technology, and X120 line pipe manufacturing technology ( Reports have been made on Patent Documents 1 and 2). However, these high-strength line pipes are also required to have brittle fracture crack propagation stopping characteristics and high-speed ductile fracture propagation stopping characteristics, and if these issues cannot be solved, they can be put into practical use as line pipes even if steel plates and steel pipes can be manufactured. It is impossible to do.

The brittle fracture crack propagation stop property needs to stop brittle fracture even if brittle fracture occurs especially from the circumferential weld that connects the line pipes. The crack propagation speed of brittle fracture will be 350 m / s or more, and brittle fracture may be a long-distance fracture ranging from 100 m to several km. Yes. As a small test for evaluating this brittle fracture crack propagation stop characteristic, it is required to have a ductile fracture surface ratio of 85% or more at a specified temperature in DWTT (Drop Weight Tear Test).

On the other hand, the high-speed ductile fracture crack propagation stopping characteristic is a phenomenon in which ductile fracture propagates at a high speed of 100 m / s or more for a long distance in the tube axis direction of the steel pipe. This high-speed ductile fracture also has a possibility of long-distance fracture ranging from 100 m to several km, and is regarded as important because of the magnitude of damage assumed. This high-speed ductile fracture has been correlated with the Charpy energy of the steel pipe, and has been prevented by securing this Charpy absorbed energy.

However, these prevention standards have been established for steel pipes with a strength level of 70 ksi (= 490 MPa) or less. For steel plates with a tensile strength of 80 ksi (= 560 MPa) or more that have been developed recently, the above parameters are not acceptable. There are concerns that it will be sufficient. No means has been established for predicting the high-speed ductile fracture propagation stopping characteristics of steel sheets having 80 ksi or more. On the other hand, for high-strength line pipes, the propagation energy of fracture due to DWTT, the crack opening angle (CTOA), or the propagation energy due to DWTT after the occurrence of ductile fracture once by pre-crack is high. The idea of dealing with ductile fracture crack propagation stopping properties has been proposed.

In order to improve the brittle crack propagation stop characteristics and ductile crack propagation stop characteristics by this DWTT, it is necessary to set the ductile / brittle transition temperature below the specified temperature. In order to lower the ductile / brittle transition temperature, that is, to improve the low temperature toughness, it is necessary to make the crystal grain size fine. The microstructure of the high-strength line pipe is a bainite / martensite-based structure. As a method for refining crystal grains in a bainite / martensite-based structure, it is known to reduce the pancake thickness. However, there is a limit to reducing the pancake thickness. Further, in the case of a bainite-martensite-based structure, it is known that {100} accumulates on a surface inclined by 40 ° with respect to the rolling surface with the rolling direction as an axis (hereinafter referred to as a 40 ° surface). {100} is a cleavage plane of iron, and if there is a brittle part such as center segregation, brittle fracture occurs from the brittle part, and brittle fracture propagates all at once to the 40 ° plane where {100} is accumulated. Therefore, it becomes difficult to shift to ductile fracture. The above is a big problem that the DWTT ductility / brittle fracture temperature in the structure mainly composed of bainite / martensite does not shift to the low temperature side. For this reason, a multiphase structure in which ferrite is generated is formed from a structure mainly composed of bainite and martensite, and a structure in which {100} is not accumulated on a 40 ° plane is created. The tissue control which suppresses was performed (patent document 3). When creating such a ferrite, the higher the strength, the more limited the amount of ferrite. When the amount of ferrite is limited, accumulation of {100} on the 40 ° plane is not suppressed, so that a brittle crack easily propagates to that plane. Another problem is to uniformly disperse ferrite throughout the steel pipe.

International Publication No. 96/023083 International Publication No. 96/023909 JP 2008-013800 A

Conventionally, as a method of refining crystal grains in a structure mainly composed of bainite and martensite, it is known to reduce the pancake thickness, but since there is an upper limit on the thickness of the slab, the pancake thickness should be reduced. Has its limits. Further, in the case of a bainite-martensite-based structure, it is known that {100} accumulates on a surface inclined by 40 ° with respect to the rolling surface with the rolling direction as an axis (hereinafter referred to as a 40 ° surface). {100} is a cleavage plane of iron, and if there is a brittle part such as center segregation, brittle fracture occurs from the brittle part, and brittle fracture propagates all at once to the 40 ° plane where {100} is accumulated. Therefore, it was a big problem that did not shift to ductile fracture.

The present invention has been made in view of such circumstances, and low temperature toughness of steel pipes used for line pipes having a structure mainly composed of bainite and martensite, particularly brittle fracture crack propagation stopping characteristics and high-speed ductile fracture. It is an object to improve the crack propagation stop characteristic.

The present inventors conducted earnest research on the conditions that should be satisfied by steel materials for obtaining high-strength linepipe steel sheets and high-strength linepipe steel pipes that have excellent tensile strength of 600 MPa or more and low-temperature toughness. It came to invent the steel plate for line pipes and the steel pipe for high-strength line pipes. And even in the structure mainly composed of bainite and martensite, the embrittlement phase such as central segregation is remarkably relaxed, and if the low temperature toughness of the place is improved, the ductile / brittle transition temperature such as DWTT can be lowered. I found it. The gist of the present invention is as follows.
(1) In mass%,
C: 0.03-0.08%,
Si: 0.01 to 0.5%,
Mn: 1.6 to 2.3%
Nb: 0.001 to 0.05%,
N: 0.0010 to 0.0050%,
Ca: 0.0001 to 0.0050%
Including
P: 0.015% or less,
S: 0.0020% or less,
Ti: 0.030% or less,
Al: 0.030% or less,
O: limited to 0.0035% or less, with the balance being Fe and inevitable impurity elements,
S / Ca <0.5
Satisfied,
Maximum Mn segregation degree: 2.0 or less,
Nb segregation degree: 4.0 or less,
Ti segregation degree: limited to 4.0 or less,
Furthermore,
A high-strength line pipe excellent in low-temperature toughness characterized by having a {100} accumulation degree limited to 4.0 or less and a tensile strength of 600 MPa or more with the rolling direction as the axis and tilted by 40 ° on the rolling surface Steel plate.
(2) Furthermore, in mass%,
Ni: 0.01 to 2.0%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 1.0%,
Mo: 0.01 to 0.60%,
W: 0.01 to 1.0%,
V: 0.01 to 0.10%
Zr: 0.0001 to 0.050%,
Ta: 0.0001 to 0.050%,
B: 0.0001 to 0.0020%
The steel sheet for high-strength line pipe excellent in low-temperature toughness according to (1), further comprising one or more of the above.
(3) By mass% REM: 0.0001 to 0.01%,
Mg: 0.0001 to 0.01%,
Y: 0.0001 to 0.005%,
Hf: 0.0001 to 0.005%,
Re: 0.0001 to 0.005%
The steel plate for high-strength line pipe excellent in low-temperature toughness according to (1) or (2), further comprising one or more of them.
(4) The steel sheet for high-strength line pipe excellent in low-temperature toughness according to any one of (1) to (3) having a bainite + martensite structure.
(5) The steel sheet for high-strength line pipe excellent in low-temperature toughness according to any one of (1) to (4), wherein the average prior austenite has an average grain size of 10 μm or less and a bainite + martensite structure.
(6) The steel sheet for high-strength line pipe excellent in low-temperature toughness according to any one of (1) to (6), which has a bainite + martensite structure and a ferrite fraction of less than 10%.
(7) The steel sheet for a high-strength line pipe excellent in low-temperature toughness according to any one of (1) to (6), wherein the maximum hardness of the center segregation part is 400 Hv or less.
(8) The base material is mass%,
C: 0.03-0.08%,
Si: 0.01 to 0.5%,
Mn: 1.6 to 2.3%
Nb: 0.001 to 0.05%,
N: 0.0010 to 0.0050%,
Ca: 0.0001 to 0.0050%
Including
P: 0.015% or less,
S: 0.002% or less,
Ti: 0.001 to 0.030%,
Al: 0.030% or less,
O: limited to 0.0035% or less, with the balance being Fe and inevitable impurity elements,
S / Ca <0.5
Satisfied,
Maximum Mn segregation degree: 2.0 or less,
Nb segregation degree: 4.0 or less,
Ti segregation degree: limited to 4.0 or less,
Furthermore,
A high-strength line pipe excellent in low-temperature toughness characterized by having a {100} accumulation degree limited to 4.0 or less and a tensile strength of 600 MPa or more with the rolling direction as the axis and tilted by 40 ° on the rolling surface Steel pipe.
(9) Furthermore, in mass%,
Ni: 0.01 to 2.0%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 1.0%,
Mo: 0.01 to 0.60%,
W: 0.01 to 1.0%,
V: 0.01 to 0.10%
Zr: 0.0001 to 0.050%,
Ta: 0.0001 to 0.050%,
B: 0.0001 to 0.0020%
The steel pipe for high-strength line pipe excellent in low-temperature toughness according to (8), characterized by containing one or more of the above.
(10) Further, by mass% REM: 0.0001 to 0.01%,
Mg: 0.0001 to 0.01%,
Y: 0.0001 to 0.005%,
Hf: 0.0001 to 0.005%,
Re: 0.0001 to 0.005%
The steel pipe for high-strength linepipe excellent in low-temperature toughness according to (8) or (9), characterized by containing one or more of them.
(11) The steel pipe for high-strength line pipe excellent in low-temperature toughness according to any one of (8) to (10), which has a bainite + martensite structure.
(12) The steel pipe for a high-strength line pipe excellent in low-temperature toughness according to (8) to (11), wherein the average prior austenite has an average particle size of 10 μm or less and has a bainite + martensite structure.
(13) The steel pipe for high-strength line pipe excellent in low-temperature toughness according to (8) to (12) having a ferrite fraction of less than 10% in a bainite-martensite structure.
(14) The steel pipe for a high-strength line pipe excellent in low-temperature toughness according to any one of (8) to (13), wherein the maximum hardness of the central segregation part is 400 Hv or less.

According to the present invention, the segregation degree of Mn, Nb, and Ti is reduced, the increase in the maximum hardness of the central segregation part is suppressed, and it is possible to produce a steel plate for line pipe and a steel pipe for line pipe excellent in low temperature toughness. For example, the industrial contribution is very remarkable.

The influence of the maximum M segregation degree on the DWTT ductile fracture surface ratio in the 0.06C-1.9Mn-Ni-Cu-Cr system is shown. The influence of the Nb segregation degree on the DWTT ductile fracture surface ratio in the 0.06C-1.9Mn-Ni-Cu-Cr system is shown. The influence of the Ti segregation degree on the DWTT ductile fracture surface ratio in the 0.06C-1.9Mn-Ni-Cu-Cr system is shown.

Hereinafter, the contents of the present invention will be described in detail. The present invention relates to an ultra-high-strength line pipe excellent in low-temperature toughness having a tensile strength (TS) of 600 MPa or more. Since the ultra-high-strength line pipe of this strength level can withstand a higher pressure than the conventional mainstream X65, it is possible to transport many gases with the same size. In the case of X65, in order to increase the pressure, it is necessary to increase the wall thickness, which increases the material cost, transportation cost, and local welding construction cost, and the pipeline laying cost significantly increases. This is the reason why an ultra-high strength line pipe having a tensile strength (TS) of 600 MPa or more and excellent in high-speed ductile fracture characteristics is required. On the other hand, when the strength is increased, it becomes difficult to manufacture a steel pipe. In this case, in order to obtain the target characteristics including seam welds, it is particularly necessary to improve the high speed fracture characteristics, improve the low temperature toughness of the base metal, improve the low temperature toughness of the weld metal and the weld heat affected zone, It is necessary to break the tube body in a burst test.

Explain the high-speed ductile fracture characteristics of the base metal. As a result of intensive studies on the fracture toughness of the base steel sheet in order to satisfy the high-speed ductile fracture characteristics of the base metal, the present inventors have found the following.

In order to improve the brittle fracture crack propagation resistance characteristics and the high-speed ductile fracture propagation characteristics, it is necessary that the base metal has high fracture propagation stop characteristics. In order to achieve this, it is known that, for example, it is important to improve the ductile fracture surface ratio in the drop weight test (DWTT) and to improve the fracture propagation energy. In the case of high strength having a tensile strength of 600 MPa or more, the structure is basically composed of bainite or martensite. In that case, the steel sheet is cooled from a temperature of Ar3 or higher to obtain a steel plate. In this case, {100}, which is a cleavage plane of iron, is accumulated at a position of 40 degrees with respect to the rolling surface with the rolling direction as an axis. Hereinafter, in this specification, a surface positioned at 40 degrees with respect to the rolling surface with the rolling direction as an axis is referred to as a “40 ° surface”. Specifically, it has more than twice the accumulation in the random case. Hereinafter, in this specification, the ratio of {100} accumulation compared to this random case is referred to as “integration degree”.

In the case of high-strength steel, for example, if the level of center segregation is poor, brittle fracture occurs from center segregation, the brittle fracture propagates along the 40 ° plane, and the DWTT ductile fracture surface rate and propagation energy decrease significantly. End up. The inventors investigated the relationship between the maximum Mn segregation degree of the center segregation part, the Ti segregation degree, the segregation degree of Nb and the DWTT ductile fracture surface ratio and the DWTT propagation energy, the maximum Mn segregation degree of the center segregation, the Ti segregation degree, It was found that the degree of segregation of Nb greatly affects the DWTT ductile fracture surface ratio or the DWTT propagation energy.

FIGS. 1 to 3 show the influence of the maximum segregation degree of Mn, Ti and Nb on the DWTT ductile fracture ratio in the 0.06C-1.9Mn—Ni—Cu—Cr system. It was found that when the maximum Mn segregation degree, Ti segregation degree, and Nb segregation degree decrease, the DWTT ductile fracture surface ratio decreases remarkably. In particular, when the maximum Mn segregation degree was 2.0 or less, the Ti segregation degree was 4.0 or less, and the segregation degree of Nb was 4.0 or less, the DWTT ductile fracture surface ratio significantly increased. Thus, the inventors consider the reason why the DWTT ductile fracture ratio is remarkably improved by decreasing the maximum Mn segregation degree, Ti segregation degree, and Nb segregation degree as follows.

When the degree of segregation of Mn increases, the Mn concentration in that region increases. Therefore, the hardenability of the center segregation portion is increased, and the hardness is greatly increased as compared with the normal portion. When the hardness in that region is increased, the low temperature toughness, specifically the fracture occurrence characteristics, is significantly reduced. Therefore, the fracture easily occurs from the center segregation, and the brittle fracture progresses on the 40 ° plane. On the other hand, when the maximum Mn segregation degree decreases, the increase in the hardness of the center segregation part is suppressed, and the resistance value of fracture increases.

On the other hand, when the Ti segregation degree and the Nb segregation degree increase, the formation of Ti and Nb carbonitrides in the central segregation part becomes remarkable, which also significantly reduces the fracture occurrence characteristics. Conversely, when the degree of segregation of Ti and Nb decreases, the generation of Ti and Nb carbonitrides in the central segregation part is suppressed, and the fracture occurrence characteristics are improved. In addition, when the maximum Mn segregation degree increases, Ca is added to suppress the formation of MnS so that the generation of MnS does not become remarkable.

Here, in the present invention, the maximum Mn segregation degree is the maximum Mn amount of the central segregation portion relative to the average Mn amount excluding the central segregation portion of the steel plate and the steel pipe. Similarly, the Nb segregation degree and the Ti segregation degree are the average Nb amount (Ti amount) averaged at the center segregation portion relative to the average Nb amount (Ti amount) excluding the center segregation portion of the steel plate and the steel pipe.

Also, when measuring the maximum Mn segregation degree, the Mn concentration distribution of the steel sheet and the steel pipe is measured by EPMA (Electron Probe Micro Analyzer) or CMA (Computer Aided Micro Analyzer) capable of image processing the measurement result by EPMA. Similarly, regarding the Nb segregation degree and the Ti segregation degree, the Nb concentration distribution and the Ti concentration distribution are measured by EPMA or CMA, respectively.

At this time, the numerical value of the maximum Mn segregation degree varies depending on the probe diameter of EPMA (or CMA). The present inventors have found that the segregation of Mn can be properly evaluated by setting the probe diameter to 2 μm. As for the Nb segregation degree and the Ti segregation degree, it was found that the segregation can be properly evaluated by setting the probe diameter to 2 μm. In the case of Nb and Ti, the maximum value cannot be accurately determined. Therefore, the average value of the measurement data, that is, the average maximum value in the thickness direction is determined. When inclusions such as MnS, TiN, and Nb (C, N) are present, the maximum Mn segregation degree, Ti segregation degree, and Nb segregation degree are apparently increased. Shall be evaluated.

Hereinafter, the reasons for limiting the chemical composition of the base material in the present invention will be described.

C: C is an element that improves the strength of steel, and as an effective lower limit, addition of 0.03% or more is necessary. On the other hand, if the C content exceeds 0.08%, the formation of carbides is promoted and the low temperature toughness of the central segregation part is impaired, so the upper limit is made 0.08% or less. Moreover, in order to suppress the low temperature toughness of a normal part, weldability, and the fall of toughness, it is preferable to make the upper limit of C amount into 0.07% or less.

Si: Si is a deoxidizing element and needs to be added in an amount of 0.01% or more. On the other hand, if the Si content exceeds 0.5%, the toughness of the weld heat affected zone (HAZ) is lowered, so the upper limit is made 0.5% or less.

Mn: Mn is an element that improves strength and toughness, and needs to be added in an amount of 1.6% or more. On the other hand, if the amount of Mn exceeds 2.3%, the low temperature toughness and the HAZ toughness of the center segregation part are lowered, so the upper limit is made 2.3% or less. In order to suppress the low temperature toughness deterioration of the center segregation part, the upper limit of the Mn content is preferably set to 2.0% or less.

Nb: Nb is an element that forms carbides and nitrides and contributes to improvement in strength. In order to obtain the effect, it is necessary to add 0.001% or more of Nb. However, if Nb is added excessively, the degree of segregation of Nb increases, and the accumulation of Nb carbonitrides is invited, resulting in a decrease in HIC resistance. Therefore, in the present invention, the upper limit of the Nb amount is 0.05% or less.

N: N is an element that forms nitrides such as TiN and NbN. In order to make the austenite grain size during heating fine by using nitrides, the lower limit of the N amount is 0.0010% or more. Is necessary. However, if the N content exceeds 0.0050%, Ti and Nb carbonitrides are likely to accumulate, and the HIC resistance is impaired. Therefore, the upper limit of the N amount is set to 0.0050% or less. In addition, when toughness etc. are requested | required, in order to suppress the coarsening of TiN, it is preferable to make the upper limit of N amount 0.0035% or less.

P: P is an impurity, and if the content exceeds 0.015%, the HIC resistance is impaired, and the toughness of the HAZ decreases. Therefore, the upper limit of the P content is limited to 0.01% or less.

S: S is an element that generates MnS that extends in the rolling direction during hot rolling, and lowers the low-temperature toughness. Therefore, in the present invention, it is necessary to reduce the amount of S, and the upper limit is limited to 0.0020% or less. In order to improve toughness, the S content is preferably 0.0010% or less. The smaller the amount of S, the better. However, it is difficult to make it less than 0.0001%. From the viewpoint of production cost, the lower limit is preferably made 0.0001% or more.

Ti: Ti is an element that is usually used for grain refinement as a deoxidizer or nitride-forming element. In the present invention, an element that lowers HIC resistance and toughness by forming carbonitrides. It is. Therefore, the upper limit of the Ti content is limited to 0.030% or less. Further, if the addition is less than 0.001%, the effect of crystal grain refinement cannot be obtained, so the lower limit is made 0.001%.

Al: Al is a deoxidizing element. However, in the present invention, when the addition amount exceeds 0.030%, an accumulation cluster of Al oxide is confirmed, so the content is limited to 0.030% or less. When toughness is required, the upper limit of Al content is preferably set to 0.017% or less. Although the lower limit of the amount of Al is not particularly limited, it is preferable to add Al in an amount of 0.0005% or more in order to reduce the amount of oxygen in the molten steel.

O: O is an impurity, and the upper limit is limited to 0.0035% or less in order to suppress oxide accumulation and improve low-temperature toughness. In order to suppress the formation of oxides and improve the base material and the HAZ toughness, the upper limit value of the O amount is preferably 0.0030% or less. The optimum upper limit of the amount of O is 0.0020% or less.

Ca: Ca is an element that generates sulfide CaS, suppresses the generation of MnS that extends in the rolling direction, and contributes significantly to the improvement of low-temperature toughness. If the addition amount of Ca is less than 0.0001%, the effect cannot be obtained, so the lower limit is made 0.0001% or more. On the other hand, if the Ca content exceeds 0.0050%, oxides accumulate and the low temperature toughness is impaired, so the upper limit is made 0.0050% or less.

In the present invention, since S is fixed by adding Ca to form CaS, the ratio of S / Ca is an important index. MnS produces | generates that ratio of S / Ca is 0.5 or more, and MnS extended at the time of rolling is formed. As a result, low temperature toughness deteriorates. Accordingly, the S / Ca ratio is set to less than 0.5.

In the present invention, one or more elements among Ni, Cu, Cr, Mo, W, V, Zr, Ta, and B may be added as elements for improving strength and toughness. it can.

Ni: Ni is an element effective for improving toughness and strength, and in order to obtain the effect, addition of 0.01% or more is necessary. However, addition of 2.0% or more reduces weldability. Therefore, the upper limit is preferably set to 2.0%.

Cu: Cu is an element effective for increasing the strength without reducing toughness, but if it is less than 0.01%, there is no effect, and if it exceeds 1.0%, cracking is likely to occur during heating of the steel slab or during welding. To do. Therefore, the content is preferably 0.01 to 1.0% or less.

Cr: Cr is effective to improve the strength of the steel by precipitation strengthening. Addition of 0.01% or more is effective, but if added in a large amount, the hardenability is increased, the bainite structure is generated, and the low temperature toughness is lowered. Let Therefore, the upper limit is preferably 1.0%.

Mo: Mo is an element that improves hardenability and at the same time forms carbonitrides and improves strength. To obtain the effect, addition of 0.01% or more is preferable. On the other hand, when Mo is added in a large amount exceeding 0.60%, the cost increases, so the upper limit is preferably made 0.60% or less. Moreover, since the low temperature toughness may decrease when the strength of the steel increases, the preferable upper limit is made 0.20% or less.

W: W is an element effective for improving strength, and is preferably added in an amount of 0.01% or more, more preferably 0.05% or more. On the other hand, if W exceeding 1.0% is added, the toughness may be lowered, so the upper limit is preferably made 1.0% or less.

V: V is an element that forms carbides and nitrides and contributes to the improvement of strength. In order to obtain the effect, addition of 0.01% or more is preferable. On the other hand, addition of V exceeding 0.10% may cause a decrease in low-temperature toughness, so the upper limit is preferably made 0.10% or less.

Zr, Ta: Zr and Ta are elements that form carbides and nitrides as well as V and contribute to the improvement of strength. In order to obtain the effect, 0.0001% or more is preferably added. On the other hand, if Zr and Ta are added excessively over 0.050%, the low temperature toughness may be lowered, so the upper limit is preferably made 0.050% or less.

B: B is an element that segregates at the grain boundaries of the steel and contributes significantly to improving the hardenability. In order to obtain this effect, 0.0001% or more of B is preferably added. Further, B is an element that generates BN, lowers the solid solution N, and contributes to the improvement of the toughness of the weld heat affected zone. Therefore, addition of 0.0005% or more is more preferable. on the other hand. When B is added excessively, segregation to the grain boundary becomes excessive, and the low temperature toughness may be lowered, so the upper limit is preferably made 0.0020%.

Furthermore, in order to control inclusions such as oxides and sulfides, one or more of REM, Mg, Zr, Ta, Y, Hf, and Re may be included.

REM: REM is an element added as a deoxidizer and a desulfurizer, and 0.0001% or more is preferably added. On the other hand, if added over 0.010%, a coarse oxide is formed, which may reduce the HIC property and the toughness of the base material and HAZ. The preferable upper limit is 0.010% or less.

Mg: Mg is an element added as a deoxidizing agent and a desulfurizing agent. In particular, a fine oxide is generated and contributes to improvement of HAZ toughness. In order to obtain this effect, 0.0001% or more of Mg is preferably added, and 0.0005% or more is more preferably added. On the other hand, if Mg is added in an amount exceeding 0.010%, the oxide tends to agglomerate and coarsen, which may lead to deterioration of HIC properties and toughness of the base material and HAZ. Therefore, it is preferable that the upper limit of the amount of Mg is 0.010% or less.

Y, Hf, Re: Y, Hf, and Re are elements that, like Ca, generate sulfides, suppress the generation of MnS elongated in the rolling direction, and contribute to the improvement of HIC resistance. In order to obtain such an effect, 0.0001% or more of Y, Hf, or Re is preferably added, and more preferably 0.0005% or more. On the other hand, if the amount of Y, Hf, Re exceeds 0.0050%, the oxides increase, and if aggregation or coarsening impairs the HIC resistance, the upper limit is preferably made 0.0050% or less.

Furthermore, in the present invention, the maximum Mn segregation degree, the Nb segregation degree, and the Ti segregation degree are 2.0 or less, 4.0 or less, and 4.0 or less, respectively.

When the maximum Mn segregation degree is 2.0 or less, the increase in hardness of the center segregation part is suppressed, and the low temperature toughness of the center segregation part is improved. Further, when the Nb segregation degree is 4.0 or less, the production of accumulated Nb (C, N) is suppressed, and when the Ti segregation degree is 4.0 or less, the production of accumulated TiN is suppressed, both of which are the central segregation part. It is possible to prevent the deterioration of the low temperature toughness.

The maximum Mn segregation degree is the maximum Mn amount of the central segregation part relative to the average Mn amount excluding the central segregation part of the steel plate and steel pipe, and the Mn concentration distribution of the steel plate and steel pipe is determined by EPMA or CMA with a probe diameter of 2 μm. Can be measured and determined. The same applies to the degree of Nb segregation and the degree of Ti segregation. The Nb concentration distribution and the Ti concentration distribution were measured by EPMA or CMA with a probe diameter of 2 μm, respectively, and the average Nb excluding the central segregation part of the steel plate and steel pipe The average Nb amount averaged at the center segregation part with respect to the amount (Nb segregation degree), the average maximum Ti amount at the center segregation part with respect to the average Ti amount excluding the center segregation part of the steel plate and steel pipe (Ti segregation degree) Is to be sought.

A method for suppressing the maximum Mn segregation degree, Nb segregation degree, and Ti segregation degree will be described below.

In order to suppress segregation of Mn, Nb and Ti, light reduction at the time of final solidification in continuous casting is optimal. The light reduction at the time of final solidification is applied to eliminate the mixing of solidified and unsolidified parts due to non-uniform cooling of the casting. it can.
When the {100} accumulation degree of the “40 ° plane” exceeds 4.0, a skewed entire brittle fracture surface is observed, and the DWTT ductile fracture surface ratio of 85% is not satisfied. Was 4.0 or less.

In continuous casting, steel slabs are usually water-cooled, but cooling at the end in the width direction is fast and cooling at the center in the width direction is strengthened. Therefore, even if the steel piece is solidified at the end portion and the center portion in the width direction, solidification is delayed at a quarter portion in the width direction, and an unsolidified portion remains inside the steel piece. Therefore, in the width direction of the steel slab, the solidified part and the unsolidified part are not uniform, and for example, the shape of the interface between the solidified part and the unsolidified part may be W-shaped in the width direction. If such non-uniform solidification occurs in the width direction, segregation is promoted, hardness increases, and low temperature toughness deteriorates.

On the other hand, in continuous casting, when light reduction is performed at the time of final solidification, an unsolidified portion is pushed out and can be solidified uniformly in the width direction. Further, if light pressure is applied after uneven solidification occurs in the width direction, the unsolidified portion cannot be effectively pushed out due to the large deformation resistance of the solidified portion.

Therefore, in order to prevent such W-type solidification from occurring, it is preferable to lightly reduce the amount of reduction while controlling the amount of reduction according to the distribution in the width direction of the central solid fraction at the final solidification position of the slab. . Thereby, center segregation is suppressed also in the width direction, and the maximum Mn segregation degree, Nb segregation degree, and Ti segregation degree can be further reduced.

Steel containing the above components is made into a steel slab by continuous casting after melting in the steel making process, and the steel slab is reheated and subjected to thick plate rolling to obtain a steel plate. In this case, the steel plate is rolled by setting the reheating temperature of the steel slab to 1000 ° C. or more, setting the reduction ratio in the recrystallization temperature region to 2 or more, and reducing the reduction ratio in the non-recrystallization region to 3 or more. Further, after the rolling is completed, water cooling is performed. It is preferable that the water cooling start temperature is set at a temperature equal to or higher than the Ar3 point and the water cooling stop temperature is set to 250 to 600 ° C. If the water cooling stop temperature is less than 250 ° C., cracking may occur. If it is this temperature range, it will become a microstructure which has a bainite martensite fraction of 90% or more. Furthermore, the average prior austenite particle size can be made 10 μm or less.
The measurement method of the average prior austenite particle size conforms to the measurement method of ASTM E112. If the rolling reduction is performed with the reduction ratio in the recrystallization temperature range of less than 2 and the reduction ratio in the non-recrystallization range of less than 3, the average prior austenite grain size cannot be made 10 μm or less. When the average prior austenite grain size is 10 μm or more, the DWTT ductile fracture surface ratio of 85% cannot be satisfied. Therefore, the average prior austenite particle size was set to 10 μm or less.
The recrystallization temperature range is a temperature range where recrystallization occurs after rolling, and is generally over 900 ° C. for the steel components of the present invention. On the other hand, the non-recrystallization temperature range is a temperature range in which recrystallization and ferrite transformation do not occur after rolling, and is generally 750 to 900 ° C. for the components of the steel of the present invention. Rolling in the recrystallization temperature range is called recrystallization rolling or rough rolling, and rolling in the non-recrystallization temperature range is called non-recrystallization rolling or finish rolling.

After the non-recrystallization rolling, the maximum hardness of center segregation can be reduced to 400 Hv or less by starting water cooling from a temperature of Ar 3 ° C. or higher and setting the water cooling stop temperature to 250 ° C. or higher. Further, when the water cooling stop temperature is set to 400 ° C. or higher, similarly, the hard martensite after transformation is partially decomposed, and the hardness can be suppressed to 350 Hv or lower. Further, if the water-cooling stop temperature is too high, the strength is lowered, so that it is necessary to add a large amount of alloy, and therefore 600 ° C. or less is preferable. In the hardness measurement method, the Mn concentration distribution of the steel plate and the steel pipe was measured by EPMA or CMA with the center segregation portion having a probe diameter of 2 μm, and the measurement location was hit in a grid pattern with a load of 25 g at a pitch of 0.5 mm. The maximum load is shown.

Next, the present invention will be described in further detail with reference to examples. Steel having chemical components shown in Table 1 was melted, and a steel piece having a thickness of 240 mm was obtained by continuous casting. In continuous casting, light reduction during final solidification was performed. The obtained steel slab was heated to 1050 to 1250 ° C., hot-rolled in a recrystallization temperature range exceeding 900 ° C., and subsequently hot-rolled in a non-recrystallization temperature range of 750 to 900 ° C. After hot rolling, water cooling was started at 700 ° C. or higher, and water cooling was stopped at a temperature of 250 to 500 ° C. Thereby, the microstructure of the steel slab obtained a structure in which the total fraction of bainite and martensite was 90% or more.

Further, the steel plate was formed into a tubular shape by C press, U press, and O press, end surfaces were tack welded, main welding was performed from the inner and outer surfaces, and then expanded to obtain a steel pipe. In addition, submerged arc welding was adopted for this welding.

Tensile test pieces, DWTT pieces, and macro test pieces were collected from the obtained steel plates and steel pipes and subjected to respective tests. DWTT was performed according to API5L3. Moreover, the segregation degree of Mn, Nb, and Ti was measured by EPMA using a macro test piece. The segregation degree was measured by EPMA with a probe diameter of 2 μm and a measurement area of total thickness × 20 mm width. The Vickers hardness of center segregation was measured according to JIS Z 2244. Vickers hardness was measured at a site where the load was 25 g and the Mn concentration was highest in the distribution of Mn concentration in the thickness direction measured by EPMA.

Table 2 shows the thickness of the steel sheet, the maximum Mn segregation degree, the Nb segregation degree, the Ti segregation degree, the maximum hardness of the central segregation part, the tensile strength, and the ductile fracture surface ratio determined by DWTT. Table 3 shows the thickness of the steel pipe, the heat input amount of the main welding, and the ductile fracture surface ratio determined by DWTT. The maximum Mn segregation degree, the Nb segregation degree, the Ti segregation degree, and the maximum hardness of the central segregation part of the steel pipe are the same as those of the steel sheet, and the tensile strength of the steel pipe is about 10% larger than that of the steel sheet.

Steels 1 to 22 and 32 are examples of the present invention, and these steel sheets have a maximum Mn segregation degree of 2.0 or less, an Nb segregation degree of 4.0 or less, a Ti segregation degree of 4.0 or less, and a central segregation portion. The maximum hardness is 400 Hv or less, and the DWTT ductile fracture surface ratios all satisfy 85% or more. The same applies to steel pipes made of these steel plates.

On the other hand, Steels 23 to 31 and 33 to 35 represent comparative examples that are outside the scope of the present invention. That is, either the basic component or the selected element is out of the scope of the present invention, or because S / Ca is 0.5 or more, the ductile fracture surface ratio by DWTT is less than 85%. I understand that.
In Steel 33, the {100} accumulation degree on the 40 ° surface exceeds 4.0, and the ductile fracture surface ratio is less than 85%. Steel 34 has a basic component element outside the scope of the present invention, and the degree of {100} accumulation on the 40 ° plane exceeds 4.0, so the ductile fracture surface ratio is less than 85%. Steel 35 has a segregation degree of Nb and a segregation degree of Ti exceeding 4.0, and a degree of {100} accumulation on the 40 ° surface exceeds 4.0, so that the ductile fracture surface ratio is 85%. Is below.

限定 Limited to the chemical components and production method of the present invention, the maximum hardness of the center segregation part and the length of the uncompressed part are limited. This effect makes it possible to manufacture steel plates for line pipes and steel pipes for line pipes that are excellent in low temperature toughness. As a result, the safety for the line pipe is greatly improved and the industrial applicability is high.

Claims

% By mass
C: 0.03-0.08%,
Si: 0.01 to 0.5%,
Mn: 1.6 to 2.3%
Nb: 0.001 to 0.05%,
N: 0.0010 to 0.0050%,
Ca: 0.0001 to 0.0050%
Including
P: 0.015% or less,
S: 0.0020% or less,
Ti: 0.001 to 0.030%,
Al: 0.030% or less,
O: limited to 0.0035% or less, with the balance being Fe and inevitable impurity elements,
S / Ca <0.5
Satisfied,
Maximum Mn segregation degree: 2.0 or less,
Nb segregation degree: 4.0 or less,
Ti segregation degree: limited to 4.0 or less,
Furthermore,
A high-strength line pipe excellent in low-temperature toughness characterized by having a {100} accumulation degree limited to 4.0 or less and a tensile strength of 600 MPa or more with the rolling direction as the axis and tilted by 40 ° on the rolling surface Steel plate.
Furthermore, in mass%,
Ni: 0.01 to 2.0%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 1.0%,
Mo: 0.01 to 0.60%,
W: 0.01 to 1.0%,
V: 0.01 to 0.10%
Zr: 0.0001 to 0.050%,
Ta: 0.0001 to 0.050%,
B: 0.0001 to 0.0020%
The steel plate for high-strength line pipe excellent in low-temperature toughness according to claim 1, comprising one or more of the following.
Further, REM: 0.0001 to 0.01% by mass%,
Mg: 0.0001 to 0.01%,
Y: 0.0001 to 0.005%,
Hf: 0.0001 to 0.005%,
Re: 0.0001 to 0.005%
The steel plate for high-strength line pipe excellent in low temperature toughness according to claim 1 or 2, characterized by containing one or more of them.
The steel sheet for high-strength line pipe excellent in low temperature toughness according to claims 1 to 3, which has a bainite + martensite structure.
The steel sheet for high-strength line pipe excellent in low-temperature toughness according to any one of claims 1 to 4, wherein the average prior austenite has an average grain size of 10 µm or less and a structure having a bainite + martensite structure.
The steel sheet for high-strength line pipe excellent in low temperature toughness according to any one of claims 1 to 5, having a bainite + martensite structure and a ferrite fraction of less than 10%.
The steel sheet for high-strength line pipe excellent in low-temperature toughness according to any one of claims 1 to 6, wherein the maximum hardness of the center segregation part is 400 Hv or less.
The base material is mass%.
C: 0.03-0.08%,
Si: 0.01 to 0.5%,
Mn: 1.6 to 2.3%
Nb: 0.001 to 0.05%,
N: 0.0010 to 0.0050%,
Ca: 0.0001 to 0.0050%
Including
P: 0.015% or less,
S: 0.002% or less,
Ti: 0.030% or less,
Al: 0.030% or less,
O: limited to 0.0035% or less, with the balance being Fe and inevitable impurity elements,
S / Ca <0.5
Satisfied,
Maximum Mn segregation degree: 2.0 or less,
Nb segregation degree: 4.0 or less,
Ti segregation degree: limited to 4.0 or less,
Furthermore,
With the rolling direction as the axis, the accumulation degree of {100} at a location inclined by 40 ° to the rolling surface is limited to 4.0 or less,
A steel pipe for high-strength line pipe excellent in low-temperature toughness characterized by having a tensile strength of 600 MPa or more.
Furthermore, in mass%,
Ni: 0.01 to 2.0%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 1.0%,
Mo: 0.01 to 0.60%,
W: 0.01 to 1.0%,
V: 0.01 to 0.10%
Zr: 0.0001 to 0.050%,
Ta: 0.0001 to 0.050%,
B: 0.0001 to 0.0020%
The steel pipe for high-strength line pipe excellent in low-temperature toughness according to claim 8, comprising one or more of the following.
Further, REM: 0.0001 to 0.01% by mass%,
Mg: 0.0001 to 0.01%,
Y: 0.0001 to 0.005%,
Hf: 0.0001 to 0.005%,
Re: 0.0001 to 0.005%
The steel pipe for high-strength line pipe excellent in low temperature toughness according to claim 8 or 9, characterized by containing one or more of them.
The steel pipe for high-strength line pipe excellent in low-temperature toughness according to any one of claims 8 to 10, which has a bainite + martensite structure.
The steel pipe for high-strength linepipe excellent in low-temperature toughness according to any one of claims 8 to 11, wherein the average prior austenite has an average grain size of 10 µm or less and has a bainite + martensite structure.
The steel pipe for high-strength line pipe excellent in low-temperature toughness according to any one of claims 8 to 12, having a ferrite fraction of less than 10% in a bainite martensite structure.
14. The steel pipe for high-strength line pipe excellent in low-temperature toughness according to any one of claims 8 to 13, wherein the maximum hardness of the center segregation part is 400 Hv or less.