JPH10265893A - 950 n/mm2 class tempered high tensile strength steel plate excellent in homogeneity in plate thickness direction and small in anisotropy of toughness and its production - Google Patents

Abstract

PROBLEM TO BE SOLVED: To produce a tempered high tensile strength steel plate excellent in homogeneity in the plate thickness direction and small in the anisotropy of toughness by composing it of the one having a specified componental compsn. contg. C, Si, Mn, Cu, Ni or the like, and the balance Fe, and in which carbon equivalent, the index of weld crack sensitivity, the quantity of alloy elements to enter into solid solutions or the like. SOLUTION: This 950 N/mm<2> class tempered high tensile strength steel plate having a compsn. contg., by mass, 0.07 to 0.18% C, 0.05 to 0.30% Si, 0.40 to 1.30% Mn, 0.01 to 0.50% Cu, 0.8 to 3.5% Ni, 0.10 to 1.20% Cr, 0.05 to 0.80% Mo and 0.005 to 0.1005 V, contg. Al, Nb and B, in which the contents of N, P and S are regulated, and the balance Fe, in which carbon equivalent Ceq (formula I) is regulated to 0.52 to 0.61%, the index of weld crack sensitivity PcM (formula II) is regulated to 0.25 to 0.31%, the solid solution quantity (formula III) of alloy elements is regulated to >=0.88%, the average grain size of Nb compounds is regulated to <=8 nm, the volume fractional rate is regulated to 1.5×10<-4> to 5.0×10<-4> , the average size No. of old austenitic grains is regulated to >=8, and the degree of elongation AI1 is regulated to <=1.5 is produced.

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a plate stock having uniformity in the thickness direction and small anisotropy of toughness, and particularly to a plate stock suitable for use as a penstock (hydraulic pipe) of a pumped storage power plant. The present invention relates to a tempered high-strength steel sheet having a thickness of 100 mm or less and a method for producing the same.

[0002]

2. Description of the Related Art In recent years, the penstock of a pumped storage power plant has tended to have a high head and a large diameter in order to respond to an increase in power demand. The steel material used for the penstock tends to have a higher tensile strength to improve the efficiency of the penstock construction and reduce costs. For this reason, there is a tendency to use a 950 N / mm ² grade steel sheet having higher tensile strength than conventional 590 N / mm ² and 780 N / mm ² grade steel sheets.

[0003] If you try the strength of the steel to 950N / mm ² class,
Usually, it is necessary to add a large amount of alloying elements such as Mn, Ni, Cr, and Mo. However, if these alloying elements are added in a large amount, the susceptibility to weld cracking increases. Therefore, instead of this method, various methods for producing Nb-added steel aiming at reducing the carbon equivalent by utilizing precipitation strengthening and the like have been conventionally proposed as the following types.

The first type is a type disclosed in Japanese Patent Application Laid-Open No. Sho 60-21847, in which Nb-added steel is subjected to normal (normal) rolling, then reheat-quenched and tempered.

[0005] The second type is a type disclosed in Japanese Patent Application Laid-Open No. 6-158160, in which Nb-added steel is subjected to controlled rolling at a low temperature, followed by reheating quenching and tempering.

A third type is a type disclosed in Japanese Patent Application Laid-Open No. 2-77521 or the like, in which Nb-added steel is controlled-rolled at a low temperature, then directly quenched and tempered.

[0007] A fourth type is an Nb-added steel as disclosed in Japanese Patent Application Laid-Open No. 2-133521 or the like, in which controlled rolling is performed at a low temperature, followed by accelerated cooling in which cooling is stopped at 600 ° C or lower, and after reheating and quenching. Tempering type.

[0008] The fifth type is a method of using Nb-added steel disclosed in claim 1 of Japanese Patent Application Laid-Open No. 2-141528 in a normal (ordinary) manner.
After re-rolling, it is reheated and quenched twice and tempered (but water cooled).

The sixth type is disclosed in Japanese Patent Laid-Open Publication No.
As an alternative to claim 1 of No. 28, a method disclosed in claim 2 of the same publication is a type in which Nb-added steel is directly quenched after controlled rolling at a low temperature, and then reheat-quenched and then tempered (but water-cooled).

[0010]

However, these conventional techniques have problems, and cannot be put into practical use as a steel material for penstocks (hydraulic iron pipes). That is, first, after the Nb-added steel is subjected to ordinary rolling, reheating and quenching-tempering is performed. In the first type, the quenching depth in the thickness direction is necessarily shallow in the reheating and quenching. Even if it satisfies, it is difficult to satisfy 950 N / mm ² at the center of the sheet thickness, the toughness is low, and the homogeneity in the sheet thickness direction is poor.

The second type is, for example, 650 to 900
Controlled rolling by applying a reduction of 30% or more at a low temperature of
The purpose of this study is to reduce the size of crystal grains by a synergistic effect with b. However, in the non-recrystallized region (less than about 930 ° C), the former austenite crystal grains become expanded processed grains even if a reduction is applied.
Since further refinement does not proceed, the effect of improving the toughness is small, and the controlled rolling is hardly effective in the central part of the sheet thickness, so that the toughness is low.

In the third type, the Nb-added steel is subjected to controlled rolling at a low temperature so that the old austenite grains are formed into expanded grains, and a mixed structure of processed martensite and lower bainite is used to combine strength and toughness. Is what you do.
However, as disclosed in Ironmaking Research No. 322, p. 99 (1986) `` Development of 100 kgf / mm Class ² High Strength Steel Sheet '', the rolling main direction (L direction) by rolling in the non-recrystallization zone rolling Due to the anisotropy of the structure in the direction perpendicular to the direction (T direction), the change in the material becomes large. In particular, the difference in toughness between the L direction and the T direction is large, and the toughness in the T direction is as low as around 100 J in the surface layer or the center of the thickness. Therefore, it is inferior in brittle crack propagation stopping characteristics in the T direction and joint toughness in the heat affected zone.

[0013] In the fourth type, during reheating and quenching subsequent to accelerated cooling, if the structure before reheating is martensite or bainite, the α to γ during reheating is reduced. Since the number of γ nucleation sites in the reverse transformation is larger than that in the case of the ferrite + pearlite structure, at the time of grain growth after γ-formation, they collide with each other and inhibit the growth, and are eventually obtained. This is based on the fact that the particle size becomes smaller. Therefore, in the state before the reheating, the main purpose is to suppress the formation of ferrite and to reverse-transform with a strained structure to make γ grains finer. However, in accelerated cooling in which the cooling stop temperature is set to 600 ° C. or lower, Nb does not form a solid solution, and the effect of suppressing the austenite grain coarsening due to precipitation of the Nb compound in the subsequent heating process for reverse transformation is poor. , Γ grains are insufficiently refined, and the effect of improving the toughness at the center of the sheet thickness is small. Also, the strength at the center of the sheet thickness is low as in the case of the re-quenched-tempered (RQ-T) material, and the strength and toughness are poor in homogeneity in the sheet thickness direction.

The fifth type has a thickness of more than 100 mm.
0N / mm Grade ² steel sheet, Ni is over 3.5 to 4.5%, and in order to realize fine γ grains during quenching, the addition of Nb suppresses the growth of γ grains during heating of steel, In addition, the main purpose is to once perform a quenching treatment to reduce the size of the structure before processing (structure before gamma formation). However, the steelmaking research No. 322, p. 99 (1986) `` 10
0kgf / mm Development of Class ² High Tensile Steel Sheets ”, reheating and quenching twice
Compared with tempering, although the toughness is improved, the strength sharply decreases, causing an increase in the amount of alloy added to secure the strength,
Manufacturing costs are increased and weld cracking susceptibility is also increased.

[0015] The sixth type is, like the fifth type, intended for a 950 N / mm ² grade steel sheet exceeding 100 mm in thickness.
Over 3.5% to 4.5%, and to achieve fine γ grains during quenching, apply a cumulative reduction of 30% or more in a temperature range of 900 ° C or more to make BN or Nb (CN) a solid solution state. 800, rolling
It finishes at ℃ or more, and the subsequent direct quenching replaces the first quenching. However, 800-900 ° C
In the low-temperature rolling control, solid solution B precipitates and hardenability decreases, and the same temperature range applies to the precipitation nose of Nb (CN) .Therefore, the amount of Nb dissolved by direct quenching is very small. It is difficult to positively utilize the pinning effect of Nb in the structure before quenching after reheating and quenching. Therefore, the difference in strength and toughness at the position in the thickness direction is large as before. In addition, water cooling after tempering causes residual stress to be inherent, distorting the steel sheet, making it difficult to form into a penstock or the like, and deteriorating dimensional accuracy.

Incidentally, the central part of the steel plate in the thickness direction is less affected by heat treatment and processing and the cooling rate is slower than the surface side part, so that the crystal grains become large and the material is most deteriorated. This is a portion that is easy to be formed, and the thicker and larger the area of the steel material, the lower the toughness. In order to provide excellent fracture toughness for a penstock on which a large impact load acts, high toughness is ensured in all thickness directions from the surface of the steel material to t / 2, and the anisotropy of toughness is small. It is necessary to improve the toughness especially at the center of the sheet thickness.

As described above, 950 N / mm ² according to the prior art is used.
Grade high-strength steel sheet is inferior in homogeneity in the thickness direction and has large anisotropy of toughness, so it can be used for penstock, which is a large-scale welded structure requiring high fracture toughness. Was not.

[0018] Accordingly, the present invention has been made in view of these prior art problems, is excellent in homogeneity in the thickness direction, also 950 N / mm ² class tempered high tensile steel with improved anisotropy of toughness and a method of manufacturing The purpose is to provide.

[0019]

The present inventors have means for solving problems] as the performance should be provided thick 950 N / mm ² class high strength steel sheet used in the penstock, etc., as shown in Table 1, in particular, of the steel sheet L direction, in order to ensure the high fracture toughness for stopping brittle crack propagation in the T direction both L direction Zen'itaAtsu direction position and the steel plate, the T direction both the 950 N / mm ² or more tensile strength, vT _rs: −
The target was to have a high toughness of 60 ° C. or less and vE _-60 : 150 J or more.

[0020]

[Table 1]

Therefore, in order to obtain the target material performance,
As a result of intensive studies on the effect of thermomechanical treatment on the solid solution amount of alloying elements and the precipitation state of Nb compounds (Nb carbonitride, etc.) on the prior austenite grain size, the following findings were obtained.

That is, after the material to be rolled is heated to a temperature equal to or higher than the solution temperature, after rolling in the recrystallization region, direct quenching (DQ) is carried out, so that C, Mo, V, Nb, B, etc. To stop martensitic transformation in the state of supersaturated solid solution to strengthen the matrix,
Nb completely solid-dissolved by direct quenching precipitates finely as an Nb compound in the subsequent temperature raising process for reheating quenching (RQ), and synergizes with the refining of the reverse transformation itself, thereby forming austenite crystal grains. By suppressing the coarsening, the DQ of FIG.
As shown in -Q material, the structure before quenching has a finer austenitic structure than other RQ materials and DQ materials (FIG. 7).
, The DQ-Q material was No. 1 and RQ in Tables 7 to 13 described later.
The No. 4 material corresponds to No. 15 and the DQ material corresponds to No. 15).
As shown in the DQ-Q material of FIG. 8, it is larger and finer than the other RQ materials (in FIG. 8, the DQ-Q material is No. 1 in Tables 7 to 13 described later, and the RQ material is No. 1). .4),
DQ-Q material is fine Nb that is coherently precipitated when RQ temperature rises.
As shown in the DQ-QT material of FIG. 9, the matrix strengthens the compound by increasing the hardness HV as compared with other RQ-T materials (in FIG. 9, the DQ-QT material is described later). Nos. 1 and RQ-T materials in Tables 7 to 13 correspond to No. 4), reheating and quenching (RQ) after direct quenching (DQ)
The DQ-Q material has higher hardenability than the RQ material because the solid solution amount of the alloy element before is significantly larger than immediately before the air cooling-RQ after rolling.

By applying these findings, in particular, the anisotropy of the prior austenite structure in the L and T directions of the steel sheet is small, the quenching depth is deep, and fine A structure mainly composed of tempered martensite having prior austenite grains was obtained. The high-strength steel sheet having this microstructure improves the strength and toughness in the central part of the sheet thickness, and also reduces the hardenability by refining the surface layer, thereby ensuring high toughness. in addition,
The present inventors have found that the steel sheet has a material having extremely low anisotropy in the L direction and the T direction, and have completed the present invention.

In the present invention, based on the above technical concept, C: 0.07 to 0.18%, Si: 0.05 to 0.30%, M
n: 0.40 to 1.30%, Cu: 0.01 to 0.50%, Ni: 0.8 to
3.5%, Cr: 0.10 to 1.20%, Mo: 0.05 to 0.80%,
V: 0.005 to 0.100%, Al: 0.005 to 0.100%, N
b: 0.008 to 0.025%, B: 0.0003 to 0.0030%, N: 0.0060% or less, P: 0.010% or less, S: 0.
005% or less, and the balance consists of Fe and unavoidable impurities, and the carbon equivalent Ceq is 0.52 to 0.61% [
Ceq = C + Si / 24 + Mn / 6 + Ni / 40 + Cr / 5 +
Mo / 4 + V / 14 (%)], the weld crack susceptibility index PcM
0.25 to 0.31% [However, PcM = C + Si / 30 + Mn / 20 + C
u / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 ×
B (%)] and the solid solution amount of the alloying element [however, the solid solution amount of the alloying element = sol.Cr + sol.Mo + sol.V + sol.B (%)]
It has a chemical composition of 0.88% or more, and the Nb compound in the steel has an average particle size of 8 nm or less and a volume fraction of 1.5 × 10 ^{-4 to} 5.
0 × 10 present in the range of ^-4 the average particle size of prior austenite grains of the steel No. 8 or more and TenShindo AI _l (where, ASTM E11
2-1995, 16.3.5, p. 237) is less than or equal to 1.5.
950N / mm Class ² tempered high-strength steel sheet with low anisotropy of toughness.

Further, as additional selective elements, Ca, T
One or more of i may be contained.

In the present invention, further, as a production method, a steel having the chemical composition is heated to 1100 ° C. or more and hot-rolled, and reduced in a temperature range of 1000 ° C. or less at a cumulative rolling reduction of 40% or more. In addition, after finishing at 930 ° C or higher, directly quenching is performed at a cooling rate at the center of the plate thickness of 3.5 ° C / sec or more to 200 ° C or less, and then re-heated to a temperature range of 870 to 950 ° C and water-quenched. And then _{tempered at} a temperature not _higher than the _Ac1 point and air-cooled. And, by this manufacturing method, more preferably, the solid solution amount of the alloy element of the steel after tempering and air cooling.
[However, the solid solution amount of the alloy element = sol.Cr + sol.Mo + sol.
V + sol.B (%)] is 0.88% or more, the Nb compound in the steel has an average particle size of 8 nm or less, and the volume fraction is 1.5 × 10 ^{-4 to} 5.
0 × 10 is present in the range of ^-4 in the average particle size of prior austenite grains of the steel No. 8 or higher and TenShindo AI _l (where, ASTM E
112 -1995,16.3.5,237 page measurements according to the method) was 1.5 or less, excellent homogeneity in the thickness direction of a plate thickness of 100mm or less, toughness small anisotropy 950 N / mm ² class tempered high of Enables production of high-strength steel sheets.

[0027]

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS First, the reasons for limiting the solid solution amount of an alloying element, the state of fine precipitation of an Nb compound and the prior austenite grain size, which are the gist of the present invention, will be described.

That is, the present inventors investigated in detail the effect of the solid solution amount of the alloy element on the strength and the effect of the precipitation state of the Nb compound and the prior austenite grain size on the toughness. The investigation was carried out under the conditions shown in Table 2. Using the steel containing 0.013% Nb of No. A in Table 3 in Examples 3 to be described later, Table 5,
From the thermomechanical treatment shown in 6, high-temperature controlled rolling (OR)-direct quenching, reheating quenching, tempering (DQ-QT) method, high-temperature controlled rolling (OR)-reheating quenching, tempering (RQ-
T) method, high temperature controlled rolling (OR)-accelerated cooling (AcC-Q)
-T, cooling stopped at 500 ° C) method, low temperature controlled rolling (CR)-
Direct quenching, reheating quenching, tempering (DQ-QT) and the like were selected.

Test specimens were taken from these steels, and the undissolved amount of alloying elements by the extraction residual analysis, the precipitation state of the Nb compound by the extraction replica method, and the prior austenite particle size by ASTM E112-95 were measured. went. In addition, as a mechanical test, tensile test at t / 4, T / 2 and surface 7mm
Below, V notch Charpy test at t / 4, T / 2
Direction and T direction, and the strength and toughness were evaluated. Table 2 also shows these results.

As for the results in Table 2, first, after OR-DQ-QT, sol.Cr + sol.Mo + so
FIG. 1 shows the relationship between the solid solution amount of the alloy element consisting of l.V + sol.B and the strength rearranged. The numbers in the figure correspond to the numbers in Table 2. From the figure, the tensile strength is DQ
-950 N / mm ²
In order to ensure the above strength, the solid solution amount of the alloying element should be 0.
It turns out that 88% or more is necessary.

Similarly, FIG. 2 shows the relationship between the cooling rate of direct quenching after rolling and the solid solution amount of alloying elements. From the figure, when the cooling rate of direct quenching is increased, D
The solid solution amount of the alloy element after Q increases, and this tendency is
It can be seen that it is taken over even after QT. Therefore, in order to obtain a solid solution amount of the alloy element after DQ-QT of 0.88% or more, it is effective to set the cooling rate during DQ to 3.5 ° C./sec or more.

Next, FIG. 3 shows the relationship between the prior austenite grain size after DQ-QT and the toughness value. As shown in the figure, as the average grain size No. of the prior austenite increases, the toughness improves, and after DQ-QT, high toughness of vT _rs : -60 ° C or less and vE _-60 : 150 J or more is obtained. It is clear that the average grain size No. of old austenite should be 8 or more.

Next, FIG. 4 shows that NQ after DQ-QT
The relationship between the precipitation state of the compound b and the prior austenite grain size No. is shown. From the figure, the grain size No.
Means that the smaller the Nb compound, the larger the average particle size
In order to make No. 8 or more, the average particle size of the Nb compound
It can be seen that it is better to set it to 8 nm or less. In order to obtain this fine Nb compound, the cooling rate during DQ is set to 3.5 ° C./se.
It is effective to make c or more.

Next, FIG. 5 shows the amount of Nb as an Nb compound precipitated in the process of raising the RQ temperature in DQ-Q,
The relationship with the average grain size No. of the former austenite is shown below.
From the figure, it can be seen that DQ
It can be seen that, in the RQ following the above, when a large amount of the Nb compound is precipitated at the time of the temperature rise, the prior austenite grains can be made fine. Accordingly, entanglement during the reverse transformation by reheating to finely precipitate the Nb compound is an extremely effective means for refining crystal grains. Further, the Nb compound finely coherently deposited has an effect of strengthening the matrix as shown in FIG.

From the above findings, the solid solution amount of the alloying element and the average grain size of the Nb compound to satisfy a tensile strength of 950 N / mm ² or more and vE _-60 : 150 J or more at all thickness direction positions. When the relationship with the diameter is determined, as is clear from FIG. 6, the chemical composition and the working heat treatment conditions are controlled so that the solid solution amount of the alloy element is 0.88% or more and the average particle diameter of the Nb compound is 8 nm or less. This is the point of the present invention.

The comparative example No. 5 in Table 2 was obtained by RQT of the sample which was cooled at 500 ° C. by accelerated cooling after high-temperature controlled rolling, while the comparative example No. 6 was obtained by controlled rolling at 850 ° C. rear,
RQT is the result of DQ. In both cases, the solid solution amount of the alloy element is small and the average particle size of the Nb compound is large, so that the strength at the center of the plate thickness is low, and the toughness of the surface layer portion and the center of the plate thickness is low. Therefore, it can be understood that only the steel sheet satisfying the requirements specified in the present invention can have the uniformity in the thickness direction and the performance with small anisotropy by adopting the high-temperature controlled rolling-DQ-QT method of 930 ° C. or more. .

[0037]

[Table 2]

The reasons for limiting the chemical components in the present invention will be described below. C is an element necessary for securing the strength as a high-tensile steel sheet. If the content is less than 0.07%, it is difficult to obtain a tensile strength of 950 N / mm ² or more, while if it exceeds 0.18%, the weld crack resistance is deteriorated. Therefore, the C content is in the range of 0.07 to 0.18%.

Although Si is an essential element for deoxidizing molten steel, if its content is less than 0.05%, its effect is insufficient. If its content is more than 0.30%, island-like martensite is formed in a welded joint, and Decreases joint toughness. Therefore, the Si content is in the range of 0.05 to 0.30%.

Mn is an indispensable element for improving the hardenability of steel and ensuring the strength. When the content is less than 0.40%, such an effect is not obtained. On the other hand, when the content is more than 1.30%, toughness and weldability deteriorate. Therefore, the Mn content is in the range of 0.40 to 1.30%.

Cu is effective in increasing the strength of the steel sheet.
In order to obtain the effect, the content needs to be 0.01% or more. However, if the content exceeds 0.50%, hot workability and weldability are deteriorated. From the above, the Cu content is in the range of 0.01 to 0.50%.

Ni has an effect of securing hardenability and improving low-temperature toughness. When the content is less than 0.8%, the tensile strength is 950 N / m2.
m ² or more cannot be secured. However, on the other hand, if the content exceeds 3.5%, the effect of improving the strength and toughness corresponding to the cost increase cannot be obtained. Therefore, the Ni content is in the range of 0.8 to 3.5%.

Cr is an element effective for increasing the strength.
If the content is less than 0.10%, this effect is not sufficiently exhibited, and if it exceeds 1.20%, the weldability is deteriorated. Therefore, the Cr content is in the range of 0.10 to 1.20%.

Mo is an element effective in increasing the strength and preventing temper softening. However, if the content is less than 0.05%, these effects cannot be sufficiently obtained. If the content exceeds 0.80%, the weldability deteriorates. Cost increases. Therefore,
The Mo content is in the range of 0.05 to 0.80%.

V is an element effective for increasing the strength and preventing temper softening. However, if the content is less than 0.005%, these effects cannot be sufficiently obtained, and if the content exceeds 0.100%,
Deterioration of toughness and weldability. Therefore, the V content is 0.005
Within the range of 0.100%.

Al is a deoxidizing element in molten steel. If its content is less than 0.005%, its effect is small, and 0.100%
If it is contained in excess of, the toughness and weldability deteriorate. Therefore, the Al content is in the range of 0.005 to 0.100%.

Since Nb is finely precipitated and exists in the austenite ground, the growth of austenite grains is suppressed by its pinning effect, the austenite grains are refined, and the strength is increased by coherent precipitation in the matrix. When the content is less than 0.008%, the effect of grain refinement is small, and when the content exceeds 0.025%, the effect of increasing the strength is saturated and the base material toughness is reduced. Therefore, the Nb content is in the range of 0.008 to 0.025%.

B is effective for increasing the hardenability and increasing the strength by adding a small amount of B, but in order to obtain such an effect, it is necessary to contain 0.0003% or more. On the other hand, 0.00
Even if the content exceeds 30%, no improvement in effect can be expected. Therefore, the content of B is set in the range of 0.0003 to 0.0030%.

In the present invention, in addition, Ca and Ti are added to
It can be selectively added depending on the target strength, toughness and processing level. Ca has a spheroidizing effect on nonmetallic inclusions and is effective in improving bending workability and toughness.
If the content is less than 0.0010%, such effects are poor,
Conversely, if the content exceeds 0.010%, improvement of the spheroidizing effect cannot be expected, and excess Ca becomes an inclusion and deteriorates toughness. Therefore, the content of Ca is 0.0010 to 0.01.
The range is 0%.

Ti is effective for miniaturization through suppression of austenite grain growth. In order to achieve such effects, 0.005% or more is required, while 0.025%
When the content exceeds the range, the base material toughness is reduced. Therefore, the Ti content is in the range of 0.005 to 0.025%.

Next, regarding the regulation of impurities, in the present invention, P, S and N are regulated respectively. First, P impairs toughness and weldability and causes hot cracking during welding. Therefore, P must be regulated to 0.010% or less.

S forms MnS inclusions with Mn, and the MnS inclusions are expanded by rolling to deteriorate bending workability and toughness. Therefore, S is 0.005%
It is necessary to regulate as follows.

N combines with B during rolling to precipitate BN, thereby reducing the effect of improving the hardenability of B. Therefore,
It is necessary to regulate to 0.0060% or less.

Further, in the present invention, the carbon equivalent Ceq.
0.61% [However, Ceq. = C + Si / 24 + Mn / 6 + Ni /
40 + Cr / 5 + Mo / 4 + V / 14 (%)].
If Ceq is less than 0.52%, strength cannot be secured, and 0.61%
If it exceeds, the weld crack resistance is poor.

Further, the welding crack susceptibility index PcM is set to 0.25 to 0.5.
31% [However, PcM = C + Si / 30 + Mn / 20 + Cu / 20 +
Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B (%)]
Range. If the PcM is less than 0.25%, the weld is softened, and if it exceeds 0.31%, the weld crack resistance is poor.

Further, in order to ensure the uniformity of the strength of the thick material in the thickness direction, as described with reference to FIG.
The solid solution amount of the alloy element represented by sol.Mo + sol.V + sol.B is set to 0.88% or more. When the solid solution amount of the alloy element is less than 0.88%, the target strength is not satisfied at the center of the plate thickness.

Next, in order to secure the material, the microstructure of the steel is specified in the present invention. First, the desired high toughness (vT _rs : -60 ° C or lower,
And vE _-60 : 150 J or more), as described in FIG.
Must be at least 8.

In order to make the old austenite grains fine, as described with reference to FIG. 4, after the Nb is completely solid-dissolved by direct quenching, the Nb compound is made fine in the subsequent temperature rise step of reheating and quenching. It is necessary to prevent the austenite grains from coarsening by reverse transformation after precipitation.
Therefore, in order to make the average particle size No. 8 or more, Nb
The compound has an average particle size of 8 nm or less and a volume fraction of 1.5
It is present in the range of × 10 ^{-4 to} 5.0 × 10 ^-4 . If the Nb compound does not satisfy any one of these conditions, the desired austenite grain size cannot be obtained, and the surface layer portion is excessively quenched, resulting in low toughness. As a result, the homogeneity of toughness is impaired in the entire thickness direction.

[0059] Furthermore, in the present invention, the austenite grain of elongation rate AI _l (However ASTM E112 -1995, Specimens with None
quiaxed Grain Shapes 16.3.5, page 237). When TenShindo AI _l is more than 1.5, the anisotropy of toughness increases, the primary rolling direction perpendicular to the direction (T direction) absorbed energy becomes lower ones. The elongation rate AI _l is, ASTM E112 -1995, Specime
ns with Nonequiaxed Grain Shapes 16.3.5, page 237, defined as follows.

[0060]

(Equation 1)

Next, the manufacturing conditions of the present invention will be described. The present invention has the object of reducing the anisotropy of toughness and having homogeneity in the thickness direction, particularly, ensuring the strength at the center of the thickness and re-establishing the surface layer and the center of the thickness. The purpose is to secure fine old austenite grains with crystal grains. Therefore, for DQ as pre-processing,
Cr, Mo, V, Nb, B, etc. are dissolved in supersaturation to ensure quenchability even in the center of the plate thickness. On the other hand, RQ is added after DQ, and Nb The point of the invention is to achieve further finer graining by entanglement of fine precipitation of the compound.

For this reason, first, regarding the slab rolling heating temperature, it is important to completely dissolve the alloy element compound (precipitate) in the slab stage at the time of rolling. However, this is an essential condition for direct quenching. If the temperature is lower than 1100 ° C., the alloy element compound cannot be completely dissolved.

Next, controlled rolling in a recrystallization region of 1000 ° C. or less is very effective for isotropic refining of austenite grains by repeating rolling and recrystallization. However, if the cumulative draft is less than 40, grain refinement cannot be sufficiently achieved. Therefore, rolling is performed at a cumulative rolling reduction of 40% or more in a temperature range of 930 to 1000 ° C. as high-temperature controlled rolling.

When the rolling finishing temperature is less than 930 ° C., solid solution elements such as Cr, Mo, V, Nb, and B precipitate in the form of a compound, and the hardenability during direct quenching is reduced. Therefore, the solid solution amount of the alloy element, which is a requirement, cannot be secured. Therefore, for the rolling finish temperature, 930
C or higher.

As a condition for direct quenching, if the cooling rate at the center of the sheet thickness is less than 3.5 ° C./sec, a complete martensitic structure is not obtained, and diffusion of C and N occurs.
Carbides and nitrides such as o, V, Nb, and B precipitate, and the solid solution amount of the alloy element, which is a requirement of the present invention, cannot be secured.
Therefore, in order to completely quench the central part of the plate thickness, the cooling rate of the central part is set to 3.5 ° C./sec or more.

Further, the cooling stop temperature of the direct quenching needs to be lower than the Mf point, and is set to 200 ° C. or lower in order to completely quench up to the center of the sheet thickness.

The reheating quenching aims at refining the martensite structure by direct quenching to reverse transformation to an austenite structure, thereby obtaining a fine-grained martensitic transformed structure by quenching from fine-grained austenite. To do so, set the quenching temperature to Ac ₃ points or more, and
The lower limit was set to 870 ° C. in order to secure the target strength, taking into account the improvement in hardenability by the method. The quenching temperature is 95
If the temperature exceeds 0 ° C., austenite grains become coarse and the low-temperature toughness of steel is impaired. Therefore, the temperature at the time of reheating should be 870-
The temperature was 950 ° C.

The tempering treatment removes the strain introduced by quenching, and finely precipitates cementite.
This is performed to improve the strength-toughness balance, but the tempering temperature is set to a temperature equal to or lower than the _Ac1 point in order to avoid an excessive decrease in strength. As a cooling method after tempering, air cooling was used in order to avoid introduction of new residual stress in consideration of bending of pen stock or the like.

[0069]

【Example】

Example 1 Next, according to the present invention described above is excellent in homogeneity in the thickness direction, a small 950 N / mm ² class tempered high tensile steel and its manufacturing method for the anisotropy of toughness, implementation This will be described with an example.

Steel slabs having the chemical components shown in Tables 3 and 4 were smelted, respectively, under the conditions shown in Tables 5 and 6 to a thickness of 50 to 100 m.
After rolling to a steel sheet of m, heat treatment such as direct quenching was performed. Then, test specimens were taken from these steel sheets, the amount of undissolved alloy elements by extraction residual analysis, and Nb by extraction replica method.
Each of the precipitation state of the compound and the prior austenite particle size was measured according to ASTM E112-95. In addition, as a mechanical test, a tensile test at t / 4 and T / 2 and a surface 7 mm below, t
The V-notch Charpy test at / 4 and T / 2 was performed in both the L direction and the T direction of the steel sheet to evaluate the strength and toughness. In addition, in order to evaluate the weldability, JIS Z 3158 Y-type welding crack test (heat input 17 kJ / cm) was used to check for cracks.
The minimum preheating temperature for crack prevention was determined. Table 7 shows the results.
To 13 are shown. In addition, the steel plates of Nos. 1 to 28 in Tables 7 to 13 are:
No. 29 is 75 mm thick, No. 30 to 32 is thick
100 mm.

In Tables 7 to 13, the steel sheet of Comparative Example No. 14 was prepared by adding low-temperature controlled rolling-DQ to the Nb-added steel in the composition of the present invention.
Although -T was applied, the strength at the center part (t / 2) of the sheet thickness did not satisfy the target strength shown in Table 1, and the toughness in the T direction was low. Further, both the strength and the toughness are inhomogeneous in the thickness direction, and the anisotropy in the L and T directions is large.

[0072] Further, the steel sheets of the comparative example of No.15, the Nb-added steel in the present invention compositions, but were subjected to hot controlled rolling -DQ-T, in Zen'itaAtsu direction, vE _-60, also poor and vT _rs, does not satisfy the target value of Table 1.

The steel sheet of Comparative Example No. 4 was prepared by subjecting the Nb-added steel in the composition of the present invention to high-temperature controlled rolling-RQ-T. 1 does not satisfy the target strength. Also, the surface layer, vE- _{60 at} t / 2, does not satisfy the target value in Table 1.

On the other hand, No. 1, 2, 8, 10, 19, 2
The steel sheets of Examples of the present invention of 0, 22 to 25, 27, and 29 to 31 were prepared by adding the Nb-added steel in the composition of the present invention to the high-temperature controlled rolling in the production conditions of the present invention.
It has been subjected to DQ-QT to have a microstructure within the scope of the present invention. In this case, the strength and toughness satisfy the target values shown in Table 1 over the entire thickness direction, have excellent homogeneity, and have almost no anisotropy in the L and T directions.
Also, the minimum preheating temperature for preventing cracking is as good as 125 ° C. or less, and satisfies the target values shown in Table 1.

On the other hand, the steel sheets of Comparative Examples Nos. 3, 5 to 7, 9, and 11 to 13 were prepared by rolling and cooling the Nb-added steel in the composition of the present invention under production conditions outside the scope of the present invention. , R-QT, the strength and toughness in the sheet thickness direction do not satisfy the target values in Table 1. Further, the homogeneity is poor in the thickness direction.

Further, the steel sheets of Comparative Examples Nos. 16 to 18, 21, 26, 28, and 32 were subjected to high-temperature controlled rolling-DQ-QT under the manufacturing conditions of the present invention for steels not having the composition of the present invention. However, none of the steel sheets has strength or toughness in the thickness direction that satisfies the target values shown in Table 1. Also, the homogeneity is poor in the thickness direction. Of these, the steel sheet of Comparative Example No. 32 had a weld cracking susceptibility index PcM exceeding the upper limit of 0.31% of the range of the present invention and was inferior in weld cracking resistance. 17kJ / cm)
° C, which does not satisfy the target value in Table 1 (125 ° C or less).

[0077]

[Table 3]

[0078]

[Table 4]

[0079]

[Table 5]

[0080]

[Table 6]

[0081]

[Table 7]

[0082]

[Table 8]

[0083]

[Table 9]

[0084]

[Table 10]

[0085]

[Table 11]

[0086]

[Table 12]

[0087]

[Table 13]

[0088]

According to the present invention, it is possible to plate thickness excellent in homogeneity in the direction, also provides a 950 N / mm ² class tempered high tensile steel and a manufacturing method thereof that improves the anisotropy of toughness . That is,
In particular, high fracture toughness for stopping brittle crack propagation in both the L direction and the T direction of the steel sheet, the position of the entire sheet thickness direction and the L
Tensile strength of 950 N / mm ² or more in both the T and T directions, and vT _rs :
It is possible to provide a thick-walled 950 N / mm class ² high-strength steel sheet having a high toughness of −60 ° C. or lower and vE _-60 : 150 J or higher,
As a result, the industrial value is great in that it allows the use of a pen stock or the like to have a high head and a large diameter.

[Brief description of the drawings]

FIG. 1 is an explanatory diagram showing the relationship between the solid solution amount of alloy elements and strength after high-temperature controlled rolling-DQ-QT in the present invention.

FIG. 2 is an explanatory diagram showing a relationship between a cooling rate during DQ after rolling and a solid solution amount of an alloy element in the present invention.

FIG. 3 is an explanatory diagram showing the relationship between the prior austenite grain size after DQ-QT and the toughness value in the present invention.

FIG. 4 is an explanatory diagram showing the relationship between the precipitation state of an Nb compound after DQ-QT and the prior austenite grain size No. in the present invention.

FIG. 5 is an explanatory diagram showing the relationship between the amount of Nb as an Nb compound precipitated upon reheating after DQ-Q and the average grain size of old austenite after DQ-Q in the present invention.

FIG. 6 is an explanatory diagram showing the relationship between the amount of Nb as a Nb compound precipitated upon reheating after DQ-Q and the average grain size of old austenite after DQ-Q in the present invention.

FIG. 7 is an explanatory diagram showing austenite crystal grains of a high temperature controlled rolling-RQ material, a DQ material, and a DQ-Q material and their sizes.

FIG. 8 is an explanatory diagram showing the precipitation behavior of Nb in a high temperature controlled rolling-RQ material, a DQ material, and a DQ-Q material.

FIG. 9 is an explanatory diagram showing the hardness of a matrix in each manufacturing process of a high temperature controlled rolling-RQ-T material and a DQ-QT material.

Claims (11)

[Claims] 1. In mass%, C: 0.07 to 0.18%, Si:
0.05 to 0.30%, Mn: 0.40 to 1.30%, Cu: 0.01 to 0.50
%, Ni: 0.8 to 3.5%, Cr: 0.10 to 1.20%, Mo:
0.05 to 0.80%, V: 0.005 to 0.100%, Al: 0.005
~ 0.100%, Nb: 0.008 ~ 0.025%, B: 0.0003 ~
0.0030%, N: 0.0060% or less, P: 0.010
% And S: 0.005% or less, respectively, the balance being Fe and unavoidable impurities, and the carbon equivalent Ceq.
0.61% [However, Ceq. = C + Si / 24 + Mn / 6 + Ni
/ 40 + Cr / 5 + Mo / 4 + V / 14 (%)] and the weld crack susceptibility index PcM is 0.25 to 0.31% [where PcM = C + Si /
30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo /
15 + V / 10 + 5 × B (%)]
[However, the solid solution amount of the alloy element = sol.Cr + sol.Mo + sol.
V + sol.B (%)] is 0.88% or more,
When the Nb compound in the steel is present in an average particle size of 8 nm or less and in a volume fraction of 1.5 × 10 ^{−4 to} 5.0 × 10 ⁻⁴ , and the average particle size No. of the prior austenite particles of the steel is 8 or more. And elongation AI
_l (according to the measurement method of ASTM E112-1995, page 16.3.5, 237) is 1.5 or less, 950 N / mm ² with excellent homogeneity in the thickness direction and small anisotropy of toughness. Grade tempered high strength steel sheet. 2. Ti: 0.005 as a selective additive element.
~ 0.025%, Ca: from among 0.0010~0.010%, 950N / mm ² class tempered high tensile steel sheet according to claim 1 containing one or two. 3. 950 N / mm ² class tempered high tensile steel sheet according to claim 1 or 2 plate thickness is less than 100mm. 4. The method according to claim 1, wherein said toughness is not higher than vT _rs -60 ° C.
_-60 : 150 J or more, 950 N / mm ² grade tempered high-strength steel sheet according to any one of claims 1 to 3. 5. applications 950 N / mm ² class tempered high tensile steel sheet according to any one any one of claims 1 to 4 is for penstock. 6. In mass%, C: 0.07 to 0.18%, Si:
0.05 to 0.30%, Mn: 0.40 to 1.30%, Cu: 0.01 to 0.50
%, Ni: 0.8 to 3.5%, Cr: 0.10 to 1.20%, Mo:
0.05 to 0.80%, V: 0.005 to 0.100%, Al: 0.005
~ 0.100%, Nb: 0.008 ~ 0.025%, B: 0.0003 ~
0.0030%, N: 0.0060% or less, P: 0.010
% And S: 0.005% or less, respectively, the balance being Fe and unavoidable impurities, and the carbon equivalent Ceq.
0.61% [However, Ceq. = C + Si / 24 + Mn / 6 + Ni
/ 40 + Cr / 5 + Mo / 4 + V / 14 (%)] and the weld crack susceptibility index PcM is 0.25 to 0.31% [where PcM = C + Si /
30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo /
15 + V / 10 + 5 × B (%)], the steel is heated to 1100 ° C. or more and hot-rolled. After finishing at ℃ or higher, the cooling rate at the center of the thickness is 3.5 ℃
Direct quenching to 200 ° C / sec or more and 200 ° C or less
Next, by reheating to water quenching to a temperature range of 870 to 950 ° C., followed characterized by air cooling tempered at a temperature below A _c1 point, the good homogeneity in the thickness direction, the toughness Manufacturing method of 950N / mm ² grade tempered high strength steel sheet with small anisotropy. 7. The steel sheet which has been tempered and air-cooled has a solid solution amount of an alloy element of the steel [however, a solid solution amount of the alloy element = sol.Cr + so.
l + Mo + sol.V + sol.B (%)] is 0.88% or more, the Nb compound in the steel has an average particle size of 8 nm or less, and the volume fraction is 1.
5 present in a range of × 10 ^-4 ~5.0 × 10 ^-4, and the average particle size of prior austenite grains of the steel No. 8 or more TenShindo AI _l (however, the ASTM E112 -1995,16.3.5,237 page ) it is 1.5 or less by the measurement method, a manufacturing method of 950 N / mm ² class tempered high tensile steel sheet according to claim 6. 8. The method according to claim 8, further comprising Ti: 0.005 as a selective additive element.
950 N / mm < ² > according to claim 6 or 7, containing one or two of Ca-0.0025% and Ca: 0.0010-0.010%.
Production method of high-grade tempered high-strength steel sheet. 9. A sheet thickness of 100 mm or less.
Manufacturing method of 950 N / mm ² class tempered high tensile steel sheet according to any one of. 10. The toughness is vT _rs : -60 ° C. or less, and
vE _-60: production method of 950 N / mm ² class tempered high tensile steel sheet according to any one of claims 6 to 9 150 J or more. 11. The use for pen stock.
Manufacturing method of 950 N / mm ² class tempered high tensile steel sheet according to to any one of 10.