JPH08311550A - Production of steel sheet for ultrahigh strength steel pipe - Google Patents

Abstract

PURPOSE: To produce a steel sheet for a steel pipe in which low temp. toughness and site weldability are simultaneously attained by subjecting a low C-Ni-Cu-Mo- Nb-Ti steel in which the content of Cu is specified to heating, executing controlled rolling and accelerated cooling and subjecting the same to forming. CONSTITUTION: The compsn. of a steel is constituted of the one contg., by weight, 0.02 to 0.10% C, <=0.6% Si, 1.0 to 2.0% Mn, <=0.015% P, <=0.0010% S, 0.3 to 1.6% Ni, 0.9 to 1.3% Cu, 0.1 to 0.5% Mo, 0.005 to 0.06% Nb, 0.005 to 0.03% Ti, <=0.06% Al, 0.001 to 0.006% N, 0.003% O, and the balance Fe. The slab having the same compsn. is rehetaed at 800 to 1000 deg.C and is rolled in such a manner that the cumulative draft at <=900 deg.C is regulated to >=70%, the cumulative rolling reduction in the two phase region of ferrite-austenite at the Ar3 to the Ar1 point is regulated to 15 to 35% and the rolling finishing temp. is regulated to 680 to 820 deg.C. It is cooled to <=400 deg.C at a cooling rate of >=10 deg.C/sec and is subjected to aging treatment at 400 to 650 deg.C.

Description

Detailed Description of the Invention

[0001]

This invention relates to the American Petroleum Institute (API)
The present invention relates to a method for producing a steel sheet having an ultrahigh strength of X100 or more (yield strength of about 689 N / mm ² or more) as standard, excellent low temperature toughness, and field weldability.

[0002]

2. Description of the Related Art Line pipes used in pipelines for long-distance transportation of crude oil and natural gas are (1) to improve transportation efficiency by increasing pressure and (2) to improve welding efficiency in the field by reducing wall thickness. There is a tendency for the tensile strength to become higher and higher. Until now, line pipes up to X80 according to the API standard have been put into practical use, but there has recently been a need for line pipes with higher strength.

At present, ultra high strength line pipes of X100 or more are manufactured by a method for manufacturing X80 class line pipes (NKK Technical Report No.
138 (1992), pp24-31 and The 7
thOffshore Mechanics and
Arctic Engineering (1988),
Volume V, pp179-185), but these line pipes have many problems in terms of low temperature toughness, field weldability, joint softening, etc.,
There is a demand for early development of an epoch-making ultra-high-strength steel pipe (line pipe) that overcomes these problems.

[0004]

DISCLOSURE OF THE INVENTION The present invention provides a technique for producing a steel plate for an ultrahigh strength steel pipe having a strength of X100 or more which can simultaneously achieve various characteristics such as low temperature toughness and field weldability.

[0005]

SUMMARY OF THE INVENTION The gist of the present invention is the weight%
And C: 0.02 to 0.10%, Si: 0.6% or less,
Mn: 1.0 to 2.0%, P: 0.015% or less, S:
0.0010% or less, Ni: 0.3 to 1.6%, Cu:
0.9-1.3%, Mo: 0.1-0.5%, Nb:
0.005-0.06%, Ti: 0.005-0.03
%, Al: 0.06% or less, N: 0.001 to 0.00
6%, O: 0.003% or less, if necessary, further C
a: 0.001 to 0.005%, V: 0.01 to 0.1
0%, Cr: 0.1 to 0.5% of one or more kinds, and the balance consisting of iron and inevitable impurities is reheated to a temperature of 800 to 1000 ° C., and then 900 ° C. or less. The cumulative reduction is 70% or more, and the cumulative reduction in the ferrite-austenite two-phase region of Ar ₃ point to Ar ₁ point is 15 to
Rolling is performed so that the rolling end temperature becomes 680 to 820 ° C. at 35%, and then 400 at a cooling rate of 10 ° C./second or more.
That is, the aging treatment is performed at a temperature of 400 to 650 ° C. after cooling to an arbitrary temperature of ℃ or less.

[0006]

The method of manufacturing the ultra high strength steel pipe of the present invention will be described in detail below. The features of the present invention are (1) 0.9-
Low C-Ni-Cu-Mo-Nb containing 1.3% Cu
After heating the Ti-based steel to the low temperature region of (2) austenite or the two-phase region of austenite-ferrite, (3)
Strictly controlled rolling in the austenite-ferrite two-phase region is followed by accelerated cooling to form a two-phase structure of fine ferrite + martensite, which provides ultra-high strength, excellent low-temperature toughness, and field weldability. Achieved at the same time.

Conventionally, Cu-precipitated steel has been used for high-tensile steel for pressure vessels (tensile strength 784 N / mm ² grade) and the like, but no development example of ultrahigh-strength line pipe of X100 or more is found. It is considered that this is because although the Cu precipitation hardened steel easily obtains strength, the low temperature toughness was insufficient as a line pipe.

First, regarding the low temperature toughness of the base material, in the pipeline, the propagation stopping characteristics as well as the brittle fracture occurrence characteristics are extremely important. In the conventional Cu precipitation hardening steel, the brittle fracture initiation characteristics represented by the Charpy characteristics were satisfactory, but the brittle fracture stopping characteristics were not sufficient. This is due to (1) insufficient microstructure miniaturization, and (2) the separation that occurs on the fracture surface of the test piece, such as the so-called Charpy impact test, has not been used (separation occurs during the impact test. It is believed that the delamination phenomenon parallel to the plate surface improves the propagation arrest property of the brittle crack by reducing the triaxial stress level at the brittle crack tip).

First, the reasons for limiting the manufacturing conditions of the present invention will be described. In the present invention, after reheating the steel slab to a temperature range of 800 to 1000 ° C, the cumulative rolling reduction of 900 ° C or less is 70%.
Above, rolling is carried out so that the cumulative rolling reduction in the ferrite-austenite two-phase region of Ar ₃ point to Ar ₁ point is 15 to 35% and the rolling end temperature is 680 to 820 ° C., and thereafter 10 ° C./sec or more. It is cooled to an arbitrary temperature of 400 ° C. or lower at a cooling rate, and is aged at a temperature of 400 to 650 ° C.

The reheating temperature of the steel slab is 800 to 1.
It is necessary to set the temperature to 000 ° C. This is because the initial austenite grains during the reheating of the steel slab are kept small and the rolling structure is refined. Furthermore, the smaller the initial austenite grains, the easier the two-phase organization of fine ferrite-martensite occurs. 1000 ° C. is the upper limit temperature at which the austenite grains at the time of reheating do not become coarse.

On the other hand, if the heating temperature is too low, the alloying elements will not be sufficiently solutionized and the desired material cannot be obtained. In addition, since heating for a long time is required to uniformly heat the billet, and further, the deformation resistance during rolling increases,
Energy cost increases, which is not preferable. Therefore, the lower limit of the reheating temperature is set to 800 ° C.

The reheated steel slab has a cumulative rolling reduction of 900 ° C. or lower of 70% or more, a cumulative rolling reduction of the ferrite-austenite two-phase region of Ar ₃ point to Ar ₁ point of 15 to 35%, and a rolling end temperature. It must be rolled to 680-820 ° C. The reason why the cumulative reduction amount at 900 ° C or less is 70% or more is to strengthen rolling in the austenite unrecrystallized region, to refine the austenite structure before transformation, and to make the structure after transformation into a ferrite-martensite two-phase structure. This is to make it an organization.

Since the X100 line pipe requires higher toughness than ever before for the sake of safety, the cumulative reduction amount must be 70% (the larger the cumulative reduction amount is, the more preferable the upper limit is not limited. ).

Further, in the present invention, the cumulative reduction amount in the ferrite-austenite two-phase region is set to 15 to 35%, and the rolling end temperature is set to 680 to 820 ° C. This is because the austenite structure finely grained in the unrecrystallized austenite region is further refined, and the ferrite is processed to strengthen the ferrite and facilitate the occurrence of separation during the impact test.

When the cumulative reduction amount in the two-phase region is 15% or less, the separation is not sufficiently generated and the propagation stopping property of the brittle crack cannot be improved. When the cumulative reduction is 35% or more, embrittlement of ferrite due to working becomes remarkable and the low temperature toughness deteriorates rather. Therefore, the range of the cumulative reduction amount in the two-phase region is set to 15 to 35%.

On the other hand, even if the cumulative reduction amount is appropriate, if the rolling temperature is inappropriate, excellent low temperature toughness cannot be achieved. If the rolling end temperature is 680 ° C. or lower, not only the effect of accelerated cooling due to the progress of ferrite transformation disappears but also embrittlement of ferrite becomes remarkable due to working, so the lower limit of the rolling end temperature was set to 680 ° C. However, the rolling end temperature is 8
At 20 ° C. or higher, refinement of the austenite structure and occurrence of separation are not sufficient, so the upper limit of the rolling end temperature was limited to 820 ° C.

After the completion of rolling, the steel sheet needs to be cooled at a cooling rate of 10 ° C./sec or more to an arbitrary temperature of 600 ° C. or less. This is transformation strengthening due to the formation of bainite structure,
This is for refining the structure and suppressing coarse Cu precipitation during cooling. If Cu precipitates during cooling, the amount of precipitation hardening after aging treatment decreases, and high strength cannot be obtained.

If the cooling rate is 10 ° C./sec or less or the water cooling stop temperature is 400 ° C. or more, improvement in strength / low temperature toughness balance due to transformation strengthening and Cu precipitation hardening cannot be expected sufficiently. The higher the cooling rate, the more effective the transformation strengthening is. The upper limit is not particularly limited, but the practically possible cooling rate is about 40 ° C./sec, although it depends on the plate thickness.

Further, the steel plate after rolling and cooling is 400 to 65
It is necessary to perform aging treatment at a temperature of 0 ° C. If it is still cooled,
Cu is hardly precipitated and Cu precipitation hardening cannot be expected. In order to increase the strength by Cu precipitation hardening (precipitation hardening by ε-Cu), aging treatment must be performed at an appropriate temperature. When the aging treatment temperature is 400 ° C. or lower, Cu precipitation is insufficient and high strength cannot be obtained, and when the aging treatment temperature is 650 ° C. or higher, Cu precipitates become coarse and the precipitation hardening ability is lost.

Next, the reasons for limiting the constituent elements will be described. The lower limit of 0.02% of C is the minimum amount for ensuring the strength and low temperature toughness of the base material and the welded portion, precipitation hardening by addition of Nb and V, and the effect of refining crystal grains. However, if the C content is too large, the low temperature toughness, the field weldability and the sour resistance are significantly deteriorated, so the upper limit was made 0.10%.

Si is an element added for deoxidation and strength improvement, but if added in a large amount, it deteriorates the field weldability and HAZ toughness, so the upper limit was made 0.6%. Deoxidation of steel is Ti
Alternatively, Al alone is sufficient, and Si does not necessarily have to be added.

Mn is an essential element for ensuring strength and low temperature toughness, and its lower limit is 1.0%, preferably 1.
3%. However, if Mn is too much, not only the hardenability of steel increases and the field weldability and NAZ toughness deteriorate,
The upper limit was set to 2.0% because it promotes center segregation of continuously cast steel pieces and also deteriorates sour resistance and low temperature toughness.

The purpose of adding Ni and Cu is to improve the strength of the low C steel of the present invention without deteriorating the low temperature toughness. It has been found that the addition of Ni and Cu rarely forms a hardened structure detrimental to the low temperature toughness in the rolled structure (especially the central segregation zone of the slab) and increases the strength, as compared with the addition of Mn, Cr and Mo. Cu sufficiently dissolves in iron even at about 800 ° C. and exhibits precipitation hardening ability to increase strength. Therefore, the Cu addition amount must be at least 0.9%.

However, if a large amount is added, local weldability and NA
Since the Z toughness is deteriorated, its upper limit is set to 1.3%. Ni is added to prevent Cu cracks during continuous casting and hot rolling, and the lower limit is 0.3%.
Is. However, if it is added in excess of 1.6%, it is not preferable for the field weldability, so the upper limit was made 1.6%.

The reason for adding Mo is to improve the hardenability of steel. Further, Mo coexists with Nb to strongly suppress recrystallization of austenite during controlled rolling, and is also effective for refining the austenite structure. In order to obtain such an effect, Mo must be at least 0.1%. However, excessive addition of Mo deteriorates HAZ toughness and field weldability, so the upper limit was made 0.5%.

In the steel of the present invention, N is an essential element.
b: 0.005%, preferably 0.01 to 0.06%,
Ti: 0.005 to 0.03% is contained. Nb contributes to refinement of crystal grains and precipitation hardening in controlled rolling, and has an action of strengthening steel. However, if Nb is added in an amount of 0.06% or more, the field weldability and HAZ toughness are adversely affected, so the upper limit was made 0.06%. When Ti is added, fine TiN is formed, and when the slab is reheated and when welding HA
It suppresses coarsening of the austenite grains of Z to make the microstructure finer and improves the low temperature toughness of the base material and HAZ.

In order to exert such an effect of TiN, it is necessary to add at least 0.005% Ti. However, if the Ti content is too large, coarsening of TiN and precipitation hardening due to TiC occur and the low temperature toughness deteriorates, so the upper limit must be limited to 0.03%.

Al is an element usually contained in steel as a deoxidizing agent and also effective in refining the structure. However, the amount of Al is 0.
If it exceeds 06%, Al-based nonmetallic inclusions increase and impair the cleanliness of the steel, so the upper limit was made 0.06%. Deoxidation is T
It is also possible to use i or Si, and it is not always necessary to add.

Further, in the present invention, P, which is an impurity element,
0.015% or less of S and O, 0.0010
% Or less and 0.003% or less. The main reason for this is to further improve the low temperature toughness of the base material and HAZ toughness. Reduction of the amount of P reduces center segregation of the continuously cast slab, prevents intergranular fracture, and improves low temperature toughness.

The reduction of the amount of S has the effect of reducing the stretched MnS and improving the sour resistance and low temperature toughness. O
The reduction of the amount reduces the oxide in the steel and is effective in improving the sour resistance and the low temperature toughness. Therefore, the lower the amount of P, S, O, the more preferable.

N forms TiN to suppress coarsening of austenite grains in the slab during reheating and in the welded HAZ to improve the low temperature toughness of the base metal and HAZ. The minimum amount required for this is 0.001%. However, if it is too large, it may cause deterioration of the HAZ toughness due to slab surface defects and solid solution N, so the upper limit must be suppressed to 0.006%.

Next, the reason for adding Ca, V and Cr will be explained. The main purpose of adding these elements to the basic composition is to increase the manufacturable plate thickness and improve the properties such as strength and toughness of the base metal without impairing the excellent characteristics of the steel of the present invention. is there. Therefore, the amount added is of a nature that should be limited by itself.

Ca controls the morphology of sulfide (MnS),
In addition to improving low temperature toughness (increasing absorbed energy in the Charpy test, etc.), it also exerts a remarkable effect in improving sour resistance. In particular, in the steel of the present invention utilizing the separation in the impact test, the absorbed energy in the Charpy test and the like tends to decrease, so that the addition of Ca is essential.

However, if the amount of Ca is less than 0.001%, there is no practical effect, and if it is added over 0.005%, it becomes C.
A large amount of aO-CaS is formed to form clusters and large inclusions, which not only impairs the cleanliness of steel, but also adversely affects on-site weldability. Therefore, the amount of Ca added is 0.001
Limited to ~ 0.005%.

V has almost the same effect as Nb, but its effect is much weaker than that of Nb. The upper limit is 0.10% in terms of field weldability and HAZ toughness.
Cr increases the strength of the base material and the welded portion, but if it is too much, it causes a remarkable deterioration in the on-site weldability and HAZ toughness. Therefore, the upper limit of the amount of Cr is 0.5%. Lower limit of V, Cr amount 0.
01% and 0.1% are the minimum amounts at which the effect on the material due to the addition of each element becomes remarkable.

[0036]

EXAMPLES Steel sheets were manufactured by various manufacturing methods from billets having various steel components by a converter-continuous casting method, and various properties were investigated. The mechanical properties were investigated in the direction perpendicular to rolling. Examples are shown in Table 1. The steel sheet produced according to the present invention has excellent strength and low temperature toughness. On the other hand, the comparative steel is not suitable in terms of chemical composition or steel plate manufacturing conditions, and either characteristic is inferior. Steel 9 has an inferior low-temperature toughness (Charpy absorbed energy, transition temperature) because it contains too much C. Steel 10 has a small amount of Mo added and an excessively large amount of Mn, so that the Charpy absorbed energy is low. Steel 11 does not have Nb added, so its strength is slightly lower than that of Nb-added steel, and its Charpy transition temperature is high (the strength / low temperature toughness balance is poor).
Has a high Charpy transition temperature because Ti is not added. Steel 13 cannot achieve the target strength because the amount of Cu added is too small.

Steel 14 has too little Ni content. as a result,
Although the mechanical properties are reasonable, many small flaws are generated on the surface of the steel pipe and it cannot be used as a line pipe. Steel 15 has an appropriate chemical composition, but has a high Charpy transition temperature because the billet reheating start temperature in the manufacturing conditions is too high. Since the reheating temperature of the steel slab is too low, the steel 16 is insufficient in solution treatment and has low strength. Steel 17 has too little cumulative reduction below 900 ° C., so low temperature toughness is a step ahead.

Steel 18 has a high Charpy transition temperature because the cumulative rolling reduction in the austenite-phyllite two-phase region is too small. Since Steel 19 has too much cumulative reduction in the two-phase region, the low temperature toughness is rather deteriorated. Steel 20 is not rolled in the two-phase region and the rolling end temperature is too high, so the low temperature toughness is poor. Steel 21 has an inferior low temperature toughness because the rolling end temperature is too low. Steel 22 has low strength because the water cooling stop temperature is too high. Steel 23 has a low strength because the aging temperature is too high. Steel 24 has a low strength because the aging temperature is too low.

[0039]

[Table 1]

[0040]

[Table 2]

[0041]

[Table 3]

[0042]

EFFECTS OF THE INVENTION According to the present invention, an ultra-high strength line pipe excellent in low temperature toughness and field weldability (API standard X100 or more)
Now, the steel plate can be manufactured stably. as a result,
The safety of the pipeline has been significantly improved, and the pipeline construction efficiency and transportation efficiency have been dramatically improved.

Claims (1)

[Claims] 1. By weight%, C: 0.02 to 0.10% Si: 0.6% or less Mn: 1.0 to 2.0% P: 0.015% or less S: 0.0010% or less Ni: 0.3-1.6% Cu: 0.9-1.3% Mo: 0.1-0.5% Nb: 0.005-0.06% Ti: 0.005-0.03% Al: 0.06% or less N: 0.001 to 0.006% O: 0.003% or less If necessary, Ca: 0.001 to 0.005
%, V: 0.01 to 0.10%, Cr: 0.1 to 0.5
%, One or two or more, and the balance consisting of iron and inevitable impurities is reheated to a temperature of 800 to 1000 ° C., and then a cumulative reduction of 900 ° C. or less is 70% or more, and an Ar ₃ point. ~ Ar ₁ point ferrite austenite 2
The cumulative rolling reduction in the phase region is 15 to 35% and the rolling end temperature is 68.
Rolling is performed at 0 to 820 ° C, and then 10 ° C /
A method for producing a steel sheet for ultrahigh-strength steel pipe, which comprises cooling to an arbitrary temperature of 400 ° C. or lower at a cooling rate of at least 2 seconds and aging at a temperature of 400 to 650 ° C.