JP2009293067A - High-tensile-strength steel material superior in formability and fatigue resistance, and manufacturing method therefor - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high-tensile-strength steel material which is suitable for an automotive structure member and is superior in formability and fatigue resistance after a cross section has been worked, and to provide a manufacturing method therefor. <P>SOLUTION: This manufacturing method includes imparting such a heat history that a cumulative heat-treatment parameter ΣAi which is defined by the expression: ΣAi=ΣäTi×(20+log ti)} (wherein ti represents heat-treatment period of time (hr) in i-th step; and Ti represents heat-treatment temperature (K) in the i-th step) satisfies 28,000 to 21,000 in a temperature range of 850 to 1,150°C and 20,000 to 13,000 in a temperature range of 500 to 700°C, and imparting such a working of which the cumulative rolling reduction in a temperature range of 850 to 1,050°C is 87 to 97%, to a base material having a composition containing 0.03 to 0.24% C and at least 0.001 to 0.15% Nb. Thereby, the steel material acquires a structure in which Nbsp/Nbsol that is a ratio of an amount of Nb in precipitates having particle sizes of smaller than 20 nm to an amount of dissolving Nb is 0.8 to 8.0. <P>COPYRIGHT: (C)2010,JPO&INPIT

Description

  The present invention relates to a high-tensile steel material suitable for automobile structural members such as a torsion beam, an ASKUL beam, a trailing arm, and a suspension arm, and particularly relates to improvement of formability and fatigue resistance. The “high-strength steel material” here refers to a steel material having a tensile strength of 590 MPa or more. The “steel material” here includes steel pipes, steel plates and the like.

  In recent years, there has been a strong demand for improving the fuel efficiency of automobiles from the viewpoint of protecting the global environment. Therefore, a thorough weight reduction of the body of an automobile or the like is aimed at. Structural members such as automobiles are no exception, and in order to achieve both weight reduction and safety, some structural members are adopting highly-strengthened ERW steel pipes. Conventionally, after forming a material (electrically welded steel pipe) into a predetermined shape, a tempering process such as a quenching process is performed to increase the strength of the member. However, adopting the tempering treatment has a problem in that the process becomes complicated, the manufacturing period of the member becomes longer, and the manufacturing cost of the member increases.

For such a problem, for example, Patent Document 1 describes a method of manufacturing an ultra-high-strength ERW steel pipe for a structural member such as an automobile. In the technique described in Patent Document 1, C, Si, Mn, P, S, Al, and N are adjusted to appropriate amounts, and B: 0.0003 to 0.003% is included, and among Mo, Ti, Nb, and V The steel material having a composition containing one or more of the above is subjected to finish rolling at 950 ° C or less at the Ar ₃ transformation point and hot rolling at 250 ° C or less to form a steel strip for pipes. This is a method for producing an ERW steel pipe, which is made into an ERW steel pipe and then subjected to an aging treatment at 500 to 650 ° C. According to this technology, it is said that an ultra-high strength steel pipe exceeding 1000 MPa can be obtained without tempering treatment by strengthening the transformation structure of B and precipitation hardening of Mo, Ti, Nb and the like.

Patent Document 2 describes a method for producing an electric-welded steel pipe having high tensile strength: 1470 N / mm ² or more and high ductility, which is suitable for automobile door impact beams and stabilizers. . In the technique described in Patent Document 2, C: 0.18 to 0.28%, Si: 0.10 to 0.50%, Mn: 0.60 to 1.80% are included, and P and S are adjusted to an appropriate range, and Ti: 0.020 to 0.050% , B: 0.0005 to 0.0050% and further subjected to normalization at 850 to 950 ° C. for an ERW steel pipe manufactured using a steel plate made of material steel having a composition containing at least one of Cr, Mo and Nb This is a method for producing an electric resistance welded steel pipe, which is further subjected to quenching treatment. According to this technique, an electric resistance welded steel pipe having a high strength of 1470 N / mm ² or more and a ductility of about 10 to 18% is obtained, which is said to be suitable for use in automobile door impact beams and stabilizers.
Japanese Patent No. 2588648 Japanese Patent No. 2814882

However, the ERW steel pipe manufactured by the technique described in Patent Document 1 is inferior in formability because of its low ductility, with an elongation El of 14% or less, and a torsion beam, axle beam, tray with press molding or hydroform molding. There is a problem that it is not suitable for automobile structural members such as ring arms and suspension arms.
On the other hand, the ERW steel pipe manufactured by the technique described in Patent Document 2 has an elongation El of at most 18% and is suitable for a stabilizer formed by bending. For the accompanying member, there is a problem that the ductility is insufficient and it is not suitable for an automobile structural member such as a torsion beam and an axle beam accompanied by press molding or hydroforming. Further, the technique described in Patent Document 2 requires a normalization process and a quenching process, has a complicated process, and has left a problem in terms of dimensional accuracy and economy.

  The present invention advantageously solves the above-described problems of the prior art, and particularly has a tensile strength of 590 MPa or more, which is suitable for automobile structural members such as a torsion beam, an axle beam, a trailing arm, and a suspension arm. It is an object of the present invention to provide a high-tensile steel material excellent in low-temperature toughness, formability, and fatigue resistance after cross-section forming processing, particularly torsional fatigue resistance, and a method for producing the same.

In the present invention, “excellent formability” means the ratio between the limit bending inner radius r and the wall thickness t obtained by a bending test based on the provisions of JIS X 2248-1996, and r / t is 3.0. The following cases shall be said.
Further, “excellent torsional fatigue resistance after cross-section forming” as used in the present invention means that the longitudinal center portion of the steel pipe is V-shaped as shown in FIG. 1 (FIG. 11 of JP-A-2001-331846). After forming the cross-section, the torsional fatigue test was performed with both ends fixed by chucking under the conditions of 1Hz and both swings, 5 × 10 ⁵ repetition fatigue limit σ _B was obtained, and 5 × 10 ⁵ repetitions obtained were obtained. The ratio between the fatigue limit σ _B and the steel pipe tensile strength TS (σ _B / TS) is 0.45 or more.

  In addition, “excellent low temperature toughness” as used in the present invention means that the longitudinal center portion of the test material (steel pipe) has a V-shaped cross section as shown in FIG. 1 (FIG. 11 of Japanese Patent Laid-Open No. 2001-331846). While still being molded, expand from the flat part of the test material so that the pipe circumferential direction (C direction) is the test piece length, and V-notch test piece (1/4 size) in accordance with JIS Z 2242 When the Charpy impact test is performed, the fracture surface transition temperature vTrs is all −50 ° C. or lower.

  The present inventors influence these properties in order to stably produce a steel material in which conflicting properties such as strength, low temperature toughness, formability, and torsional fatigue resistance after cross-section forming are compatible at a high level. A systematic study was conducted on various factors, especially the composition of steel materials and production conditions. As a result, a steel material having a composition containing C and at least one carbide forming element X, preferably Nb, is subjected to thermal history and processing under specific conditions, and the amount of fine precipitates and solid solution amount of the carbide forming element X It has been found that by controlling the ratio in a proper range, it is possible to produce a high strength steel material having a predetermined high strength, high formability, excellent low temperature toughness, and excellent torsional fatigue resistance after cross-section forming. .

The present invention has been completed based on the above-described findings and further studies. That is, the gist of the present invention is as follows.
(1) In mass%, C: 0.03 to 0.24%, including at least one carbide forming element X, X amount Xsp in precipitates having a particle size of less than 20 nm, and solid solution X amount A high-tensile steel material excellent in formability and fatigue resistance, characterized by having a structure in which the ratio of Xsol and Xsp / Xsol falls within a predetermined range.
(2) In (1), Nb in a precipitate having a composition containing C: 0.03 to 0.24% by mass%, Nb: 0.001 to 0.15% as the carbide forming element X, and having a particle size of less than 20 nm A high-strength steel material excellent in formability and fatigue resistance, characterized by having a structure in which the amount Nbsp is a ratio of the solid solution Nb amount Nbsol and the Nbsp / Nbsol is 0.80 to 8.0.
(3) In (2), the composition is mass%, C: 0.03-0.24%, Si: 0.002-0.95%, Mn: 1.01-1.99%, Al: 0.01-0.08%, Nb: 0.001-0.15% In addition, P, S, N, and O are adjusted to include P: 0.019% or less, S: 0.010% or less, N: 0.008% or less, and O: 0.003% or less, with the balance being Fe and inevitable impurities. A high-strength steel material having a composition comprising:
(4) In (2) or (3), in addition to the above-described composition, the carbide forming element X may further include, in mass%, V: 0.001 to 0.15%, Ti: 0.001 to 0.15%, Mo: 0.001 to 0.45%, W: When the composition contains one or more selected from 0.001 to 0.15% and the structure further contains V, the amount of Vsp in the precipitate having a particle size of less than 20 nm is solidified. The ratio of dissolved V amount Vsol, Vsp / Vsol is 0.20 to 2.0, and when Ti is contained, the ratio of Ti amount Tisp and solid solution Ti amount Tisol in precipitates having a particle size of less than 20 nm, Tisp / Tisol Is 0.8 to 8.0, and when Mo is contained, the ratio of Mo amount Mosp to solid solution Mo amount Mosol in the precipitate having a particle size of less than 20 nm, Mosp / Mosol is 0.10 to 1.0, and W is contained. In the case, a high-tensile steel material characterized by having a structure in which the ratio of W amount Wsp and solid solution W amount Wsol in a precipitate having a particle size of less than 20 nm, Wsp / Wsol is 0.10 to 1.0.
(5) In any one of (2) to (4), in addition to the above composition, in addition to mass, Cr: 0.001 to 0.45%, Cu: 0.001 to 0.45%, Ni: 0.001 to 0.45%, B: 0.0001 A high-tensile steel material characterized by having a composition containing one or more selected from -0.0009%.
(6) In any one of (2) to (5), in addition to the above-described composition, the high-strength steel material further includes Ca: 0.0001 to 0.005% by mass%.
(7) In any one of (2) to (6), the high-tensile steel material has an arithmetic average roughness Ra of 2 μm or less on the inner and outer surfaces of the pipe, a maximum height roughness Rz of 30 μm or less, and a ten-point average roughness Rz. _JIS is 20 μm or less, and in addition to the above structure, the structure in the region from the outermost surface and the innermost surface to 50 μm in the thickness direction has a ferrite phase with a volume ratio of 60% or more. A high-strength steel material, characterized in that it is a hollow tubular body having an average crystal grain size in a circumferential cross section of 2 to 8 µm.
(8) More than the temperature at which the carbide of carbide forming element X is dissolved in a steel material having a composition containing C: 0.03 to 0.24% and containing at least Nb: 0.001 to 0.15% as carbide forming element X. In order to obtain a steel material by applying a predetermined heat history and a predetermined processing, the predetermined heat history is expressed by the following equation (1)
ΣAi ＝ Σ {Ti ・ (20 + log ti)} (2)
(Where ti: heat treatment time (hr) in the i-th process, Ti: heat treatment temperature (K) in the i-th process)
The cumulative heat treatment parameter ΣAi defined in the above is a heat history satisfying 28000-21000 in the temperature range of 850-1150 ° C and 20000-13000 in the temperature range of 500-700 ° C, and the predetermined processing is 850-1050 ° C A method for producing a high-strength steel material having excellent formability and fatigue resistance, characterized in that the cumulative rolling reduction in the temperature range is 87 to 97%.
(9) In (8), the composition contains, in mass%, C: 0.03-0.24%, Si: 0.002-0.95%, Mn: 1.01-1.99%, Al: 0.01-0.08%, and a carbide forming element X And Nb: 0.001 to 0.15%, and further containing P, S, N, and O, adjusted to P: 0.019% or less, S: 0.010% or less, N: 0.008% or less, O: 0.003% or less, A method for producing a high-strength steel material, characterized in that the balance is composed of Fe and inevitable impurities.
(10) In (8) or (9), in addition to the above composition, the carbide-forming element X further includes, in mass%, V: 0.001 to 0.15%, Ti: 0.001 to 0.15%, Mo: 0.001 to 0.45%, W : A method for producing a high-strength steel material, characterized in that the composition contains one or more selected from 0.001 to 0.15%.
(11) In any one of (8) to (10), in addition to the above composition, in addition to mass, Cr: 0.0001 to 0.45%, Cu: 0.001 to 0.45%, Ni: 0.001 to 0.45%, B: 0.0001 A method for producing a high-strength steel material, characterized in that the composition contains one or more selected from -0.0009%.
(12) In any one of (8) to (11), in addition to the above composition, the method further includes a composition containing Ca: 0.001 to 0.005% by mass%.

  According to the present invention, a high-tensile steel material having a tensile strength of 590 MPa or more, excellent low-temperature toughness, excellent formability, and excellent torsional fatigue resistance after cross-section forming processing can be manufactured at low cost. Has an exceptional effect. Moreover, according to the present invention, there is an effect that it contributes remarkably to the improvement of the characteristics of the automobile structural member.

First, the reasons for limiting the composition of the high-strength steel material of the present invention will be described. Hereinafter, unless otherwise specified, mass% is simply expressed as%.
The high-tensile steel material of the present invention has a composition containing C: 0.03 to 0.24% and further containing at least one carbide forming element X, preferably Nb: 0.001 to 0.15%. The carbide forming element X is one or more of V, Ti, Mo, and W in addition to Nb.

C: 0.03-0.24%
C is an element that increases the strength of steel through precipitation strengthening as a precipitate (carbide) by combining with solid solution strengthening or carbide forming element X, and is an essential element for securing steel strength and fatigue strength. It is. Such an effect is recognized when the content is 0.03% or more. If the content is less than 0.03%, the desired amount of precipitates cannot be obtained, and the desired torsional fatigue resistance characteristics cannot be ensured. On the other hand, if the content exceeds 0.24%, the ductility of the steel material decreases, and the desired formability cannot be secured, and the low temperature toughness decreases. For this reason, C was limited to the range of 0.03-0.24%. In addition, Preferably it is 0.08 to 0.20%.

Nb: 0.001 to 0.15%
Nb is an element that combines with C and precipitates as a carbide, contributing to ensuring a desired high strength and improving fatigue strength. In order to obtain such an effect, a content of 0.001% or more is required. On the other hand, when the content exceeds 0.15%, the ductility decreases remarkably due to excessive precipitation of precipitates. For this reason, Nb is limited to 0.15% or less. In addition, Preferably it is 0.010 to 0.049%.

In the present invention, the carbide forming element X is selected from V: 0.001 to 0.15%, Ti: 0.001 to 0.15%, Mo: 0.001 to 0.45%, and W: 0.001 to 0.15% in addition to the above-described Nb. Moreover, you may contain 1 type (s) or 2 or more types.
V, Ti, Mo, and W are all carbide-forming elements like Nb, are precipitated as fine carbides, and have an effect of improving fatigue strength. In addition to Nb, one kind is optionally added. Alternatively, two or more kinds can be selected and contained.

V combines with C and precipitates as fine carbides, and has the effect of improving fatigue strength. Such an effect is manifested at a content of 0.001% or more. On the other hand, if it exceeds 0.15%, formability and low temperature toughness are lowered. For this reason, when it contains, it is preferable to limit V to 0.001 to 0.15% of range. More preferably, it is 0.06% or less.
In addition, Ti has an action of first combining with N to reduce the solute N, and contributes effectively to ensuring the formability of the steel material, and Ti other than nitride, that is, excess Ti, is C and Bonds and precipitates as carbides and has the effect of improving fatigue strength. Such an effect becomes remarkable when the content is 0.001% or more. However, when the content exceeds 0.15%, the strength is significantly increased by precipitates (carbides), and the ductility is remarkably decreased accordingly, and the moldability is decreased. . For this reason, when it contains, it is preferable to limit Ti to the range of 0.001 to 0.15%. In addition, More preferably, it is 0.0010 to 0.080%.

Also, Mo, like Nb, V, etc., binds to C and precipitates as carbides, and has an action of improving fatigue strength. Such an effect is manifested at a content of 0.001% or more, but a content exceeding 0.45% reduces moldability. For this reason, when it contains, it is preferable to limit Mo to 0.001 to 0.45% of range. More preferably, it is 0.12 to 0.20%.
W, like Nb, V, etc., binds to C and precipitates as carbides, and has the effect of improving fatigue strength. Such an effect is manifested at a content of 0.001% or more. On the other hand, if it exceeds 0.15%, formability and low temperature toughness are lowered. For this reason, when it contains, it is preferable to limit W to 0.001 to 0.15% of range. More preferably, it is 0.06% or less.

Other than the above-mentioned carbide-forming elements such as C and Nb, the following components are preferable.
Si: 0.002 to 0.95%
Si is a ferrite-forming element and promotes the transformation of austenite (γ) → ferrite (α) during the thermal history and contributes to improvement of formability. Such an effect becomes remarkable when the content is 0.002% or more. On the other hand, if the content exceeds 0.95%, the surface properties are lowered. For this reason, it is preferable to limit Si to 0.002 to 0.95% of range. In addition, More preferably, it is 0.10 to 0.30%.

Mn: 1.01-1.99%
Mn is an element that contributes to an increase in the strength of the steel material and has an effect of improving the fatigue strength. Such an effect is manifested at a content of 1.01% or more. If the content exceeds 1.99%, the ductility is remarkably lowered, and the desired formability cannot be ensured. For this reason, it is preferable to limit Mn to the range of 1.01-1.99%. Further, it is more preferably 1.20 to 1.80%.

Al: 0.01-0.08%
Al is an element that acts as a deoxidizer during steelmaking, and has an action of binding to N and suppressing the growth of austenite grains in the thermal history, refining crystal grains, and improving fatigue strength. Such an effect comes to be recognized when the content is 0.01% or more. On the other hand, if the content exceeds 0.08%, the effect is saturated and an effect commensurate with the content cannot be expected, which is economically disadvantageous. In addition, the oxide inclusions are increased, and the fatigue resistance is significantly deteriorated. For this reason, it is preferable to limit Al to the range of 0.01 to 0.08%. More preferably, it is 0.02 to 0.06%.

In the present invention, the impurity elements P, S, N, and O are adjusted to include P: 0.019% or less, S: 0.010% or less, N: 0.008% or less, and O: 0.003% or less. preferable.
P is an element having an adverse effect of lowering the low temperature toughness and lowering the electroweldability through solidification co-segregation with Mn, and is preferably reduced as much as possible. If the content exceeds 0.019%, the above-described adverse effects become remarkable, so P is limited to 0.019% or less.

  S exists as inclusions such as MnS in steel, and lowers formability and fatigue resistance as a starting point of fine cracks and fatigue cracks during molding. Moreover, it is an element which has the bad influence which reduces the electric resistance weldability, low temperature toughness, etc. of steel materials, and it is preferable to reduce as much as possible. If the content exceeds 0.010%, the above-described adverse effects become remarkable, so S is preferably made 0.010% as an upper limit. More preferably, it is 0.005% or less.

If N remains as solid solution N in the steel, N is an element having an adverse effect of lowering the formability and low temperature toughness of the steel material. In the present invention, N is preferably reduced as much as possible. If the content exceeds 0.008%, the above-described adverse effects become remarkable, so N is preferably made 0.008% as an upper limit. More preferably, it is 0.0049% or less.
O is an element that exists as an oxide inclusion in steel and has an adverse effect of lowering the fatigue resistance and low temperature toughness of the steel, and is preferably reduced as much as possible in the present invention. When the content exceeds 0.003%, the above-described adverse effects become remarkable, so O is preferably made 0.003% as the upper limit. More preferably, it is 0.002% or less.

In addition to the above-described components, one or more selected from Cr: 0.001 to 0.45%, Cu: 0.001 to 0.45%, Ni: 0.001 to 0.45%, B: 0.0001 to 0.0009%, and / or Or Ca: 0.001-0.005% can be contained.
Cr, Cu, Ni, and B are all elements that improve the fatigue strength or supplement the effect of improving the fatigue strength, and can be selected and contained as necessary.

Cr is an element having an effect of improving fatigue strength, and such an effect is manifested when the content is 0.001% or more. On the other hand, if the content exceeds 0.45%, the moldability is lowered. For this reason, when it contains, it is preferable to limit Cr to 0.001 to 0.45% of range. Further, it is more preferably 0.08 to 0.29%.
Both Cu and Ni are elements that complement the effect of improving fatigue strength, but are also elements that further improve the corrosion resistance of steel materials. These effects appear when the content is 0.001% or more. On the other hand, if the content exceeds 0.45%, the moldability is lowered. For this reason, when it contains, it is preferable to limit Cu to 0.001 to 0.45% and Ni to 0.001 to 0.45%. More preferably, both Cu and Ni are 0.2% or less.

Similarly, B is an element that complements the effect of improving fatigue strength. Such an effect is manifested when the content is 0.0001% or more. On the other hand, if the content exceeds 0.0009%, the moldability is lowered. For this reason, when it contains, it is preferable to limit B to 0.0001 to 0.0009% of range.
Ca has an action of controlling the form of so-called inclusions in which the expanded inclusions (MnS) are granular inclusions (Ca (Al) S (O)). It is an element which has the effect | action which suppresses the generation | occurrence | production of the fine crack and fatigue crack at the time of shaping | molding, and improves a moldability, a fatigue-resistant characteristic, and low temperature toughness through form control of such an inclusion. Such an effect becomes remarkable when the content is 0.0001% or more. However, when the content exceeds 0.005%, the non-metallic inclusions are increased, and the fatigue resistance is deteriorated. For this reason, when it contains, it is preferable to limit Ca to 0.0001 to 0.005% of range.

The balance other than the above components is Fe and inevitable impurities.
Further, as described above, the high-tensile steel material of the present invention is a solid comprising C, a composition containing at least one carbide forming element X, preferably Nb, and an X amount Xsp in a precipitate having a particle size of less than 20 nm. The ratio of dissolved X amount Xsol, the structure in which Xsp / Xsol falls within a predetermined range, and the ratio of Nb amount Nbsp and solid solution Nb amount Nbsol in precipitates having a particle size of less than 20 nm when containing Nb, Nbsp / Nbsol has a structure of 0.80 to 8.0. In addition, the quantification method according to size of the precipitate will be described later.

Fine precipitates (Nb carbides) having a particle size of carbide forming element X of less than 20 nm increase the yield strength of the steel and increase the fatigue strength. However, when fine precipitates having a particle size of less than 20 nm are excessively precipitated, the formability of the steel material is lowered, and the stress relaxation characteristics at the stress concentration portion are also lowered at the stage of initial fatigue crack generation.
On the other hand, the solid solution X (solid solution Nb) has the effect of suppressing the deterioration of the formability of the steel material, in particular the bending formability, through the refinement of the structure and increasing the fatigue strength of the steel material. However, since the contribution of solid solution X (solid solution Nb) to fatigue resistance is small compared to the contribution of fine precipitates with a particle size of less than 20 nm, the amount of fine precipitates (Nb carbide) with a particle size of less than 20 nm It is necessary to balance the amount of dissolved X (the amount of dissolved Nb).

FIG. 2 shows the relationship between r / t, (σ _B / TS) after cross-section forming, and Nbsp / Nbsol. In the present invention, the formability (bending formability) is represented by the ratio of the critical bending radius r to the wall thickness t, r / t, and the torsional fatigue resistance is 5 × 10 ⁵ repeated fatigue limit σ _B after cross-sectional forming. Of steel pipe tensile strength TS, represented by (σ _B / TS).
From FIG. 2, it can be seen that by setting Nbsp / Nbsol in the range of 0.80 to 8.0, it is possible to obtain a steel material having both desired excellent bend formability and excellent torsional fatigue resistance after cross-section forming. When Nbsp / Nbsol is less than 0.80, the torsional fatigue resistance shown by (σ _B / TS) is deteriorated, and when Nbsp / Nbsol is more than 8.0, the formability represented by r / t is lowered. Therefore, in the present invention, Nbsp / Nbsol is limited to the range of 0.80 to 8.0. In addition, Preferably it is 1.6-6.0.

Further, the carbide forming element X may further contain one or more selected from V, Ti, Mo, and W in addition to Nb.
When V is contained as the carbide forming element X, the ratio of the V amount Vsp in the precipitate having a particle size of less than 20 nm to the solid solution V amount Vsol, Vsp / Vsol is adjusted to 0.20 to 2.0. Thereby, desired excellent bending formability and excellent torsional fatigue resistance after cross-section forming can be combined. In addition, Preferably it is 0.40-1.6.

Further, when Ti is contained as the carbide forming element X, the ratio of Ti amount Tisp in precipitates having a particle diameter of less than 20 nm to solid solution Ti amount Tisol, Tisp / Tisol is adjusted to 0.80 to 8.0. Thereby, desired excellent bending formability and excellent torsional fatigue resistance after cross-section forming can be combined. In addition, Preferably it is 1.6-6.0.
Further, when Mo is contained as the carbide forming element X, the ratio of Mo amount Mosp in precipitates having a particle size of less than 20 nm to solid solution Mo amount Mosol, Mosp / Mosol is adjusted to 0.10 to 1.0. Thereby, desired excellent bending formability and excellent torsional fatigue resistance after cross-section forming can be combined. In addition, Preferably it is 0.20-0.80.

When W is contained as the carbide forming element X, the ratio of the W amount Wsp in the precipitate having a particle size of less than 20 nm to the solid solution W amount Wsol, Wsp / Wsol is adjusted to 0.10 to 1.0. Thereby, desired excellent bending formability and excellent torsional fatigue resistance after cross-section forming can be combined. In addition, Preferably it is 0.20-0.80.
Furthermore, the high-strength steel material of the present invention has the above-described composition, and has a surface roughness measured in accordance with the provisions of JIS B 0601-2001. The arithmetic average roughness Ra of the inner and outer surfaces of the pipe is 2 μm or less, and the maximum height. The roughness Rz is 30 μm or less and the ten-point average roughness Rz _JIS is 20 μm or less. In addition to the above structure, the structure in the region from the outermost surface and the innermost surface to 50 μm in the thickness direction has a volume ratio of 60. It is preferable to make a hollow tubular body (steel pipe) having a ferrite phase of at least% and having an average crystal grain size of 2 to 8 μm in the circumferential cross section of the ferrite phase.

If the volume fraction of the ferrite phase in the above region is less than 60%, the ductility is lowered and the desired formability cannot be ensured. In addition, local thinning, rough surface, and fine cracks accompanying the reduction in ductility become stress-concentrated portions, and the fatigue resistance is greatly reduced. For this reason, it is preferable to limit the volume fraction of the ferrite phase in the above region to 60% or more. More preferably, it is 75% or more.
Here, the ferrite phase includes acicular ferrite, Widmanstatten ferrite, bainitic ferrite and the like in addition to polygonal ferrite. Examples of the second phase other than the ferrite phase include carbide, pearlite, bainite, martensite, and a mixed phase thereof.

  Moreover, it is preferable that the average crystal grain diameter in the circumferential cross section of the ferrite phase in the above region is 2 to 8 μm. If the average particle size is less than 2 μm, desired moldability cannot be ensured, and local thinning, rough surface, and fine cracks become stress-concentrated portions, and the fatigue resistance is greatly reduced. On the other hand, when the average particle size exceeds 8 μm, the moldability is lowered, the surface hardness is lowered, and the fatigue resistance is lowered. For this reason, it is preferable that the average particle diameter of the ferrite phase in the region from the outermost surface of the tube and the innermost surface to 50 μm in the thickness direction is limited to a range of 2 to 8 μm. More preferably, it is 5 μm or less.

Below, the preferable manufacturing method of this invention steel material is demonstrated.
In the present invention, the steel material (slab, steel plate, steel pipe, etc.) having the above composition is heated to a temperature higher than the temperature at which the carbide of the carbide forming element X is dissolved (solid solution temperature), and then the heat history such as hot rolling. In addition, it is preferable to apply a heat history such as processing and heat treatment under appropriate conditions to obtain a high-strength steel material in which the ratio between the amount of fine precipitates and the solid solution amount of carbide-forming elements is within the appropriate range as described above.

  The steel material is first heated to a temperature equal to or higher than the temperature at which the carbide of the carbide forming element X is dissolved (solid solution temperature). If the heating temperature is less than the solid solution temperature of carbide, coarse carbide remains, and a desired amount of fine precipitates having a particle size of less than 20 nm cannot be secured, resulting in a decrease in fatigue strength. The upper limit of the heating temperature is not particularly limited, but is preferably about 1300 ° C. at which crystal grain coarsening becomes remarkable. Here, the solid solution temperature of the carbide is the equilibrium solid solution temperature of each carbide defined by the following equation.

log [Nb] [C + 12 / 14N] =-6770 / T + 2.26
log [Ti] [C] = − 7000 / T + 2.75
log [V] [C] =-9500 / T + 6.72
In addition, in the case of Mo and W, it shall be more than the A _c1 transformation point defined by the following formula ( _note that elements not contained are calculated as zero).

A _c1 transformation point (℃) = 721-10.7Mn-16.9Ni + 29.1Si + 16.9Cr + 290As + 6.38W
(Here, Mn, Ni, Si, Cr, As, W: content of each element (mass%))
Next, predetermined heat history and predetermined processing are performed so that the ratio between the amount of fine precipitates and the solid solution amount of the carbide-forming element is within the appropriate range as described above.
Predetermined thermal history is the cumulative heat treatment parameter ΣAi
ΣAi = Σ {Ti · (20 + logt _i )}
(Where t _i : heat treatment time (hr) in the i th step, Ti: heat treatment temperature (K) in the i th step)
However, it is preferable that the thermal history be such that ΣAi = 28000-21000 in the temperature range of 850-1150 ° C. and ΣAi = 20000-13000 in the temperature range of 500-700 ° C. The predetermined processing is preferably processing in which the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. is 87 to 97%. Thereby, the ratio of the amount of fine precipitates of the carbide forming element X and the solid solution amount of the carbide forming element X in the proper range as described above can be ensured. Here, the cumulative heat treatment parameter ΣAi is a well-known parameter, and is described, for example, as the Lanson-Miller parameter in “Improved 5th edition heat treatment p.164” edited by the Japan Iron and Steel Institute.

  In the austenite region of 850 ° C. to 1150 ° C., precipitates having a particle size of 20 nm or more are mainly precipitated. The amount of precipitation is related to the cumulative heat treatment parameter ΣAi including temperature and time. When ΣAi in the austenite region of 850 ° C. to 1150 ° C. exceeds 28000, the ratio between the amount of fine precipitates having a particle size of carbide forming element X of less than 20 nm and the solid solution amount of carbide forming element X in a desired range. , Can not secure. If ΣAi in the austenite region of 850 ° C. to 1150 ° C. is less than 21000, a desired cumulative rolling reduction in processing cannot be secured.

The predetermined processing is processing in which the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. is 87 to 97%.
In the above-described heat history of the austenite region, processing is applied to impart processing strain. The temperature range of 850 to 1050 ° C corresponds to the recrystallization temperature range of austenite, and the processing in this temperature range contributes to improving the formability (bending formability) by refining the structure of the steel surface layer. . When the cumulative rolling reduction is less than 87%, the structure in the region from the outermost surface and the innermost surface to the thickness direction of 50 μm has a ferrite phase with a volume ratio of 60% or more, and the circumferential cross section of the ferrite phase In this case, the desired crystal structure with an average crystal grain size of 2 to 8 μm cannot be obtained, and the moldability deteriorates. On the other hand, if the cumulative rolling reduction exceeds 97%, precipitation is promoted and the amount of precipitates increases, so that desired Nbsp / Nbsol cannot be secured. For this reason, it is preferable to limit the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. in the range of 87 to 97%. In addition, Preferably it is 90 to 95%.

  In the ferrite region of 500 to 700 ° C., precipitates having a particle size of less than 20 nm are mainly deposited. The amount of precipitation is related to the cumulative heat treatment parameter ΣAi including temperature and time. When ΣAi in the ferrite region of 500 to 700 ° C. exceeds 20000, the amount of fine precipitates having a particle size of less than 20 nm becomes excessive and formability is deteriorated. On the other hand, when ΣAi is less than 13000, the amount of fine precipitates having a particle size of less than 20 nm is small and fatigue strength is reduced. For this reason, it is preferable to limit (SIGMA) Ai in the ferrite region of 500-700 degreeC to the range of 20000-13000. More preferably, it is 19000-15000.

  Fig. 3 shows the cumulative rolling reduction in the temperature range of 850 to 1050 ° C and the cumulative heat treatment in the temperature range of 500 to 700 ° C on the Nbsp / Nbsol value in steels containing C and Nb as a carbide-forming element. The influence of the relationship with the parameter ΣAi is shown. The cumulative heat treatment parameter ΣAi in the temperature range of 850 to 1150 ° C. was about 25000. From FIG. 3, in the region where the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. is 87 to 97% and the cumulative heat treatment parameter ΣAi in the temperature range of 500 to 700 ° C. is surrounded by the range of 20000 to 13000, Nbsp / Nbsol It can be seen that the value is in the desirable range of 0.8 to 8.0.

The high-tensile steel material of the present invention is heat-treated with the thermal history of the above-described conditions, for example, a welded steel pipe, a cold-rolled steel sheet, the thermal history of the above-described conditions, or the processing of the above-described conditions, for example, heat This includes steel strips, thick plates, and seamless steel pipes.
The following is a method for manufacturing a hollow tubular body (welded steel pipe), using an electric-forged pipe process that uses a hot-rolled hot-rolled steel strip as a closed cross-section material (hollow tubular body: steel pipe) by means of electric-welding welding. An example of this will be described.

  A hot rolled steel strip produced by heating, heat history, and processing of the above-described conditions to the steel material having the above composition may remain hot rolled, but the surface is subjected to pickling, shot blasting, etc. It is preferable to remove the black skin. Furthermore, from the viewpoint of corrosion resistance and coating film adhesion, the steel strip can be subjected to surface treatment such as galvanization, aluminum plating, nickel plating, organic coating treatment, and the like. The steel strip that has been pickled or surface-treated is preferably made into a steel pipe by applying an electric sewing pipe having a width drawing ratio of 10% or less.

  Note that the material of the welded steel pipe is not limited to the hot-rolled steel strip. There is no problem even if the above-described hot-rolled steel strip is used as a raw material, and a cold-rolled annealed steel strip that has been subjected to cold rolling or annealing, or a surface-treated steel strip that has been subjected to various surface treatments is used. Also, instead of electric sewing pipes, pipe forming processes that combine roll forming, press-cut cross section of the cut plate, heat treatment such as cold, warm, hot diameter reduction and annealing after pipe forming Also good. Furthermore, there is no problem even if laser welding, arc welding, plasma welding, spot welding or the like is used instead of electric seam welding. The steel material of the present invention, for example, a hollow tubular body (closed cross-section material), exhibits excellent characteristics as it is formed into a member, but additionally has a high surface by heat treatment such as residual stress removal annealing or shot peening after forming the member. There is no problem in applying hardness and applying compressive residual stress.

In addition, the hollow tubular body manufactured in this way has a surface roughness measured in accordance with the provisions of JIS B 0601-2001. Preferably, the arithmetic average roughness Ra of the inner and outer surfaces of the pipe is 2 μm or less, the maximum A surface-like hollow tubular body having a height roughness Rz of 30 μm or less and a ten-point average roughness Rz _JIS of 20 μm or less is preferable. Thereby, especially the stress concentration resulting from the unevenness | corrugation of the internal and external surface of a pipe | tube used as the starting point of fatigue is relieve | moderated, and a torsional fatigue-proof characteristic improves notably.

  In addition to the precipitate distribution described above, the hollow tubular body produced in this way preferably further has a volume ratio of 60% in the region from the outermost surface and the innermost surface to 50 μm in the thickness direction. It has the above ferrite phase, and has a structure in which the average crystal grain size in the circumferential cross section of the ferrite phase is 2 to 8 μm. Thereby, formability and torsional fatigue resistance are significantly improved.

  The steel slab having the composition shown in Table 1 is heated to 1230 ° C, which is higher than the solid solution temperature of the carbide of the carbide forming element X contained, and processed at the cumulative reduction rate in the temperature range of 850 to 1050 ° C shown in Table 2. In addition, a thermal history in which the cumulative heat treatment parameter ΣAi in the temperature range of 850 to 1150 ° C. is about 24500 is applied, and subsequently, in the temperature range of 500 to 700 ° C., the range of ΣAi = 11500 to 22500 as shown in Table 2 A heat history was applied to obtain a hot rolled steel strip having a thickness of about 3 mm. Table 2 summarizes the heat history of each hot-rolled steel strip and the cumulative rolling reduction ratio.

Next, after pickling these hot-rolled steel strips, slitting was applied to a predetermined width dimension to obtain a pipe material. These pipe materials are continuously formed into open pipes, and the open pipes are subjected to electro-welded pipes that are electro-welded by high-frequency resistance welding, and welded steel pipes (product pipes) that are hollow tubular bodies (outer diameter 89.1 mmφ × The wall thickness was about 3 mm). In the case of an electric sewing tube, the width drawing ratio was 4%.
Test materials were collected from the obtained welded steel pipes, and subjected to tensile tests, bending tests, quantitative tests of precipitates and solid solutions, structural observations, torsional fatigue tests, low temperature toughness tests, and surface roughness tests. The test method was as follows.
(1) Tensile test From these welded steel pipes, a JIS No. 12 test piece is cut out in accordance with the provisions of JIS Z 2201 and the tensile test is carried out in accordance with the provisions of JIS Z 2241 so that the pipe axis direction is the tensile direction. Tensile properties (tensile strength TS, yield strength YS, El) were obtained.
(2) Bending test From these welded steel pipes, strip test pieces equivalent to JIS No. 1 are collected in the circumferential direction, and the pipe inner surface and pipe outer surface are on the punch side in accordance with the provisions of JIS X 2248-1996. Then, a bending test was performed, and a ratio of the limit bending radius r to the wall thickness t, r / t, was obtained. When the limit bending radius r is different between the inner surface and the outer surface of the tube, the larger one is adopted.
(3) Precipitate amount / Solid solution quantitative test From these welded steel pipes, a test piece of 20 x 30 mm x wall thickness of about 3 mm was cut out, and 10 vol% acetylacetone-1 mass% tetramethylammonium chloride-methanol electrolyte (hereinafter referred to as "electrolytic solution") , 10% AA electrolyte solution) was subjected to constant current electrolysis at a current density of 20 mA / cm ² . After electrolysis, remove the specimen with deposits on the surface from the electrolyte and immerse it in an aqueous solution of sodium hexametaphosphate (500 mg / l) (hereinafter referred to as SHMP aqueous solution) to give ultrasonic vibration. The precipitate was peeled from the specimen and extracted into an aqueous SHMP solution. Subsequently, the SHMP aqueous solution containing the precipitate was filtered using a hole diameter 100 nm filter and a hole diameter 20 nm filter in this order. The residue on the filter after filtration and the filtrate are analyzed using an ICP emission spectrophotometer, and the residues on the filter and the absolute value of carbide forming elements X (Mo, V, W, Ti, Nb) in the filtrate are analyzed. The amount of X (Mo, V, W, Ti, Nb) contained in precipitates having a particle size of more than 100 nm, precipitates having a particle size of 100 to 20 nm, and precipitates having a particle size of less than 20 nm are measured. The quantities Xlp, Xmp, Xsp were obtained. The electrolytic weight was calculated by measuring the weight of the specimen after peeling the deposit and subtracting it from the specimen weight before electrolysis.

The solid solution amount of the carbide forming element X was determined by measuring the concentration of Fe in the liquid as the carbide forming element X and the comparative element using the electrolytic solution after electrolysis as an analysis solution and using ICP mass spectrometry. Based on the obtained concentration, the concentration ratio of the carbide-forming element X to Fe is calculated, and the content of the carbide-forming element X in the solid solution state (solid solution) is further multiplied by the Fe content in the specimen. X amount) was calculated. Note that the Fe content in the specimen can be determined by subtracting the total content of components other than Fe from 100%.
(4) Microstructure observation From these welded steel pipes, specimens for microstructural observation were collected so that the observation surface had a circumferential cross-section, and the structure was polished, nital-corroded, and scanned using a scanning electron microscope (magnification: 3000 times). Were observed, imaged, and the obtained structural photograph was measured for the fraction of the ferrite phase (volume fraction) and the average grain size of the ferrite phase using an image analyzer. Note that the area occupied by each crystal grain of the ferrite phase was measured, and further converted into the equivalent circle diameter to obtain the ferrite crystal grain size (μm). The measurement position is divided into three blocks in the depth direction from the outermost surface and the inner surface to a depth of 50 μm, and the total average (n = 6) of the measured values for each block is obtained. The structure fraction of the ferrite phase and the average crystal grain size of the ferrite phase were used.
(5) Torsional fatigue test A test pipe (length: 1500 mm) was collected from the obtained welded steel pipe. Then, about 1000 mmL of the central portion of the collected test tube material is subjected to cross-section molding processing so that the circumferential cross-section is V-shaped as shown in FIG. 1 (Japanese Patent Laid-Open No. 2001-331846), and torsional fatigue It was set as the test material for a test.

The torsional fatigue test was performed by changing the stress level variously under the conditions of 1 Hz and both swings, and the number of repetitions N until the fracture at the load stress S was obtained. The resulting asked to S-N diagram than 5 × 10 ⁵ repeatedly limit σ _B (MPa), the ratio of the tube tensile strength TS for _{σ B (σ B / TS)} , were evaluated resistance torsional fatigue characteristics. The load stress S was measured by first conducting a torsion test with a dummy piece, confirming the fatigue crack position, and attaching a triaxial strain gauge at that position.
(6) Low temperature toughness test From the obtained welded steel pipe, a test pipe (length: 1500 mm) was sampled, and about 1000 mmL of the central part of the collected test pipe was shown in FIG. 1 (FIG. 11 of JP-A-2001-331846). ) As shown in Fig. 2, the cross-section is processed so that the cross-section in the circumferential direction becomes V-shaped. From the flat part of the test material, the pipe is circumferentially extended (C direction) to the test piece length. A V-notch specimen (1/4 size) was cut out in accordance with the regulations, a Charpy impact test was performed, the fracture surface transition temperature vTrs was obtained, and the low temperature toughness was evaluated.
(7) Surface roughness test The surface roughness of the inner and outer surfaces of the obtained welded steel pipe was measured using a stylus type roughness meter in accordance with the provisions of JIS B 0601-2001, As the roughness parameters, arithmetic average roughness Ra, maximum height roughness Rz, and ten-point average roughness Rz _JIS were determined. The measurement direction of the roughness curve was the circumferential direction (C direction) of the tube, the low-frequency cut-off value was 0.8 mm, and the evaluation length was 4 mm. As the representative value, the larger one of the inner surface and the outer surface was adopted.

  The obtained results are shown in Tables 3 and 4.

In any of the present invention, elongation El is 16% or more, r / t is 3.0 or less and excellent formability, (σ _B / TS) is 0.45 or more and excellent torsional fatigue resistance after cross-section forming, cross-section forming It is a welded steel pipe (steel material) that has excellent low-temperature toughness with a later fracture surface transition temperature vTrs of -50 ° C or lower. On the other hand, in the comparative example out of the scope of the present invention, the r / t is larger than 3.0 and the formability is lowered, or (σ _B / TS) is less than 0.45 and the torsional fatigue resistance after the cross-section forming process is lowered. Or vTrs exceeds -50 ° C, resulting in a product tube with low temperature toughness.

Comparative examples (pipe No. 1, No. 1) in which the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. is smaller than the preferred range of the present invention or ΣAi in the temperature range of 500 to 700 ° C. is smaller than the preferred range of the present invention. 4, No.7, No.10, No.13, No.16, No.19, No.22, No.25, No.28) are all Nb and Ti in precipitates with a particle size of less than 20 nm. , V, Mo, W amount and Nb, Ti, V, Mo, W solid solution ratio At least one of Nbsp / Nbsol, Tisp / Tisol, Vsp / Vsol, Mosp / Mosol, Wsp / Wsol The torsional fatigue resistance is reduced when the (σ _B / TS) after cross-section forming is smaller than the invention range and less than 0.45. In addition, a comparative example (pipe No. 3, in which the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. is larger than the preferred range of the present invention or ΣAi in the temperature range of 500 to 700 ° C. is larger than the preferred range of the present invention. No.6, No.9, No.12, No.15, No.18, No.21, No.24, No.27, No.30) are all Nbsp / Nbsol, Tisp / Tisol, Vsp / Vsol , Mosp / Mosol, Wsp / Wsol is larger than the range of the present invention, El is less than 16%, r / t is larger than 3.0, moldability is reduced, and (σ _B / TS) is less than 0.45 and the torsional fatigue resistance is degraded.

Further, although the cumulative rolling reduction in the temperature range of 850 to 1050 ° C. and ΣAi in the temperature range of 500 to 700 ° C. are both within the preferred range of the present invention, the comparative example (tube) whose composition is outside the preferred range of the present invention. In each of No. 31 to No. 40), the r / t is larger than 3.0 and the formability is reduced, and (σ _B / TS) after cross-section forming processing is less than 0.45, and the torsional fatigue resistance is reduced. Yes. In these comparative examples, except for the pipes No. 31 and No. 40, the vTrs exceeds −50 ° C. and the low temperature toughness is also lowered.

In the comparative examples (pipe No. 31, No. 32) in which the C amount is outside the preferred range of the present invention, Nbsp / Nbsol is larger than the present range, and the amounts of Nb, V, W, Ti, and Mo are In the comparative examples (pipe No. 36 to No. 40) exceeding the preferred range, at least one of Nbsp / Nbsol, Tisp / Tisol, Vsp / Vsol, Mosp / Mosol, Wsp / Wsol is smaller than the scope of the present invention. The r / t exceeds 3.0, the formability is reduced, and the (σ _B / TS) after the cross-section forming process is less than 0.45, and the torsional fatigue resistance is reduced. In the comparative example (pipe No. 33) in which the Si amount is less than the preferred range of the present invention, the ferrite structure fraction is lower than the preferred range of the present invention, r / t exceeds 3.0, the moldability is reduced, and the cross-section molding is performed. The (σ _B / TS) after processing is less than 0.45, and the torsional fatigue resistance is degraded. On the other hand, the comparative example (pipe No. 34) in which the Si amount exceeds the preferred range of the present invention has a large surface roughness on the inner and outer surfaces of the pipe, and the comparative example (pipe No. 35) containing no carbide forming element is ferrite. The average crystal grain size is larger than the preferred range of the present invention, r / t exceeds 3.0, the formability decreases, and (σ _B / TS) after cross-section forming processing is less than 0.45, the torsional fatigue resistance is reduced. ing.

  The surface roughness of the inner and outer surfaces of product pipes (pipe No.1 to No.33, No.35 to No.40) has an arithmetic average roughness Ra of 2 μm or less and a maximum height roughness Rz of 30 μm or less. The ten-point average roughness RzJIS was as good as 20 μm or less. In Table 3, the surface roughness of the product tubes (tubes No. 1 to No. 33) is omitted.

It is explanatory drawing which shows the cross-section shaping | molding processing state of the test material used for the torsional fatigue test and low temperature toughness test in an Example. Ratio of Nb content Nbsp and solid solution Nb content Nbsol in precipitates having a particle size of less than 20 nm, Nbsp / Nbsol, ratio of critical bending radius r and thickness t in bending test, r / t, after cross-section forming It is a graph which shows the relationship with ((sigma) _B / TS). It is a graph which shows the influence of the relationship between the cumulative reduction in the temperature range of 850 to 1050 ° C. and the cumulative heat treatment parameter ΣAi in the temperature range of 500 to 700 ° C. on the Nbsp / Nbsol value.

Claims (12)

  C: 0.03 to 0.24% by mass and having a composition containing at least one carbide-forming element X, and an X amount Xsp in a precipitate having a particle size of less than 20 nm and a solid solution X amount Xsol A high-tensile steel material excellent in formability and fatigue resistance, characterized by having a structure in which the ratio, Xsp / Xsol falls within a predetermined range.   The ratio of the Nb content Nbsp in the precipitate having a particle size of less than 20 nm and the solid solution Nb content Nbsol having a composition containing C: 0.03-0.24% and Nb: 0.001-0.15% by mass%, Nbsp / A high-strength steel material excellent in formability and fatigue resistance, characterized by having a structure in which Nbsol is 0.80 to 8.0. The composition is in weight percent,
C: 0.03-0.24%, Si: 0.002-0.95%,
Mn: 1.01-1.99%, Al: 0.01-0.08%,
Nb: 0.001 to 0.15% contained,
Further, P, S, N, and O are adjusted to include P: 0.019% or less, S: 0.010% or less, N: 0.008% or less, and O: 0.003% or less, with the balance being Fe and inevitable impurities. The high-tensile steel material according to claim 2, wherein In addition to the above composition, one or more kinds selected from V: 0.001 to 0.15%, Ti: 0.001 to 0.15%, Mo: 0.001 to 0.45%, and W: 0.001 to 0.15% in mass%. And the composition further comprises:
In the case of containing V, the ratio of the V amount Vsp in the precipitate having a particle size of less than 20 nm to the solid solution V amount Vsol, Vsp / Vsol is 0.20 to 2.0,
When Ti is contained, the ratio of the Ti amount Tisp in the precipitate having a particle size of less than 20 nm to the solid solution Ti amount Tisol, Tisp / Tisol is 0.80 to 8.0,
When Mo is contained, the ratio of Mo amount Mosp in precipitates having a particle size of less than 20 nm to solid solution Mo amount Mosol, Mosp / Mosol is 0.10 to 1.0,
3. When W is contained, the ratio of the W amount Wsp in the precipitate having a particle size of less than 20 nm to the solid solution W amount Wsol, and the structure has a structure in which Wsp / Wsol is 0.10 to 1.0. Or the high-tensile steel material of 3.   In addition to the above composition, one or more selected from Cr: 0.001-0.45%, Cu: 0.001-0.45%, Ni: 0.001-0.45%, B: 0.0001-0.0009% in mass% The high-tensile steel material according to claim 2, wherein the high-tensile steel material has a composition containing   The high-strength steel material according to any one of claims 2 to 5, wherein in addition to the composition, the composition further contains Ca: 0.0001 to 0.005% by mass. The high-strength steel material has an arithmetic average roughness Ra of the pipe inner and outer surfaces of 2 μm or less, a maximum height roughness Rz of 30 μm or less, and a ten-point average roughness Rz _JIS of 20 μm or less. The structure in the region from the outer surface and the innermost surface to the thickness direction of 50 μm has a ferrite phase with a volume ratio of 60% or more, and the average crystal grain size in the circumferential section of the ferrite phase is 2 to 8 μm The high-strength steel material according to any one of claims 2 to 6, which is a hollow tubular body. A steel material having a composition containing C: 0.03 to 0.24% and containing at least Nb: 0.001 to 0.15% as a carbide forming element X at a mass% is heated to a temperature higher than the temperature at which the carbide of the carbide forming element X is dissolved. After that, given a predetermined heat history and predetermined processing,
The predetermined heat history is a heat that satisfies a cumulative heat treatment parameter ΣAi defined by the following formula (1) of 28000 to 21000 in the temperature range of 850 to 1150 ° C and 20000 to 13000 in the temperature range of 500 to 700 ° C. History and
A method for producing a high-tensile steel material excellent in formability and fatigue resistance, wherein the predetermined processing is processing in which a cumulative rolling reduction in a temperature range of 850 to 1050 ° C is 87 to 97%.
Record
ΣAi ＝ Σ {Ti ・ (20 + log ti)} (1)
Where ti: heat treatment time in the i-th process (hr)
Ti: Heat treatment temperature in the i-th process (K) The composition is in weight percent,
C: 0.03-0.24%, Si: 0.002-0.95%,
Mn: 1.01-1.99%, Al: 0.01-0.08%
Nb: 0.001 to 0.15% as a carbide forming element X, P, S, N, O, P: 0.019% or less, S: 0.010% or less, N: 0.008% or less, O: 0.003% The method for producing a high-strength steel material according to claim 8, characterized in that the composition is adjusted as follows and the balance is composed of Fe and inevitable impurities.   In addition to the above-mentioned composition, the carbide forming element X is selected in mass% from V: 0.001 to 0.15%, Ti: 0.001 to 0.15%, Mo: 0.001 to 0.45%, W: 0.001 to 0.15%. The method for producing a high-tensile steel material according to claim 8 or 9, wherein the composition contains seeds or two or more kinds.   In addition to the above composition, one or more selected from Cr: 0.001-0.45%, Cu: 0.001-0.45%, Ni: 0.001-0.45%, B: 0.0001-0.0009% in mass% The method for producing a high-strength steel material according to any one of claims 8 to 10, characterized in that the composition contains s.   The method for producing a high-strength steel material according to any one of claims 8 to 11, wherein the composition further contains Ca: 0.0001 to 0.005% by mass% in addition to the composition.

Images