JP2000204434A - Ferritic heat resistant steel excellent in high temperature strength and its production - Google Patents

Abstract

PROBLEM TO BE SOLVED: To impart stably high creep rupture strength to the steel even in the case of use at high temp. for a long time by allowing spherical or disk- shaped coherent precipitates having specified diameter to exist in the crystal grains of steel having a specified componential compsn. at specified density, also allowing prescribed intergranular precipitates to exist and specifying the ratio between the minor axis and the major axis thereof. SOLUTION: Ferritic heat resistant steel having a compsn. contg., by weight, 0.01 to 0.25% C, <=0.03% P, <=0.015% S, 0.5 to <8% Cr, 0.05 to 0.5% V, 0 to 0.2% Nb, 0 to 0.1% Ti, 0 to 0.2% Ta and 0 to 0.1% N, in which, in the crystal grains, spherical or disk-shaped coherent precipitates of 2 to 30 nm diameter are present at the density of >=1 piece/μm3, also, one or more kinds of intergranular precipitates among cementite, M7C3 type carbides and M23C6 type carbides are present, all intergranular precipitates contain >=2 wt.% V, and the ratio between the minor axis and the major axis (minor axis/major axis) is >=0.5 is prepd. In this way, the steel usable at about 400 to 625 deg.C can be obtd.

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION The present invention relates to a cast iron and forged steel product such as a heat exchanger, a steel pipe for piping, a heat-resistant valve, a connection joint, etc. which is used in the fields of boilers, chemical industry, nuclear power and the like.
Low and medium C with excellent creep strength at high temperature of 0 ° C or higher
Related to r-ferritic heat-resistant steel.

[0002]

2. Description of the Related Art Heat resistant steels used at a high temperature of 400 ° C. or higher include austenitic stainless steel, high Cr ferrite steel having a Cr content of 9 to 12%, and low and medium Cr ferrite having a Cr content of several percent. It is roughly divided into steel and carbon steel.

[0003] The steel type depends on the use environment (temperature, pressure, etc.).
It is appropriately selected in consideration of economic efficiency.

[0004] Among these steel types, low and medium C
The r-ferritic heat-resistant steel is generally a heat-resistant steel having a tempered martensite or a tempered bainite structure containing several percent of Cr and, if necessary, alloying elements such as W, Mo, Ni, and Co.

[0005] Low- and medium-Cr ferritic heat-resistant steels are characterized by the fact that they contain Cr and thus have higher oxidation resistance than carbon steel.
It is excellent in high temperature corrosion resistance and high temperature strength. Also, compared to austenitic stainless steel, it is inferior in high-temperature strength, but is significantly cheaper, has a smaller coefficient of thermal expansion, and is less expensive than high Cr ferritic steel, and has toughness, weldability, and It is characterized by having excellent thermal conductivity.

As a typical example of such low and medium Cr ferritic steels, STBA20 (0.5 Cr
-0.5Mo), STBA22 (1.0Cr-0.5Mo), STBA23
(1.25Cr-0.5Mo), STBA24 (2.25Cr-1.0Mo), STB
A25 (5.0Cr-0.5Mo) and the like are known.

[0007] High temperature strength is extremely important in designing a pressure-resistant member, and it is desirable that the strength be high regardless of the operating temperature.
In particular, in a heat-resistant and pressure-resistant steel pipe used for boilers, chemical industries, nuclear power, and the like, the wall thickness of the pipe is determined according to the high-temperature strength of the material.

[0008] The high temperature strength of low and medium Cr ferritic steels is improved by solid solution strengthening and precipitation strengthening.

The improvement of high-temperature strength by solid solution strengthening is generally carried out by incorporating appropriate amounts of C, Cr, Mo and W. However, when used at high temperatures for a long time, coarsening of carbides and precipitation of intermetallic compounds occur. And the creep strength on the high temperature and long time side decreases. To increase the strength, it is conceivable to increase the amount of solid solution elements to increase solid solution strengthening.However, the addition beyond the solid solution limit promotes the precipitation of these elements, and on the contrary, toughness, workability, welding Deterioration of performance.

Improvement in high-temperature strength by precipitation strengthening is performed by adding precipitation strengthening elements such as V, Nb, Ti, Ta and the like. Such low and medium Cr ferrite steels are disclosed in JP-A-55-6458, JP-A-57-131349, and JP-B-61-61349.
-34501, JP-B-6-6771, JP-A-6-100929, JP-A
Many proposals have been made in JP-A-8-134584, JP-A-8-158022, etc., and they have been put to practical use.

Further, 1Cr-1Mo-0.25V steel which is a material for turbines and 2.25Cr-1Mo-Nb steel which is a structural material for fast breeder reactors are well known as precipitation strengthened low and Cr ferritic steels. I have.

However, in the case of precipitation strengthening, the following problems occur unless proper structure control is performed.

(A) An unused material or a material used for a short time has a high strength, but when exposed to a high temperature for as long as 10,000 hours or more, the precipitation effect is reduced and a stable strength cannot be obtained. This is due to carbides in unused and short-time use materials,
This is because nitrides and intermetallic compounds contribute to precipitation strengthening, but due to aging that occurs during long-term use at high temperatures, these precipitates are agglomerated and coarse, and the precipitation strengthening ability is lost.

(B) In the precipitation-strengthened steel, since the inside of the grains is strengthened, the grain boundaries become relatively weak, and the toughness and corrosion resistance deteriorate.

If the high-temperature strength of the low and medium Cr ferritic steel can be further increased, the following advantages can be obtained.

1) Conventionally, austenitic stainless steel or high Cr ferritic steel has been used to ensure high-temperature strength even in a use environment in which high-temperature corrosion resistance is not so severely required, but instead of austenitic stainless steel. If low and medium Cr ferritic steels are used, the characteristics of low and medium Cr ferritic steels, for example, excellent weldability, can be utilized.

2) In conventional applications, it is possible to reduce the wall thickness, thereby improving the thermal conductivity, improving the thermal efficiency of the plant itself, and the thermal fatigue associated with starting and stopping the plant. The load can be reduced.

3) The plant can be made compact and the manufacturing cost can be reduced by reducing the wall thickness to reduce the weight.

[0019]

The problem to be solved by the present invention is that
High creep strength at a high temperature of about 0 to 625 ° C,
Another object of the present invention is to provide a low and medium Cr ferritic steel exhibiting stable strength even when used for a long time in such a temperature range, and a method for producing the same.

[0020]

The gist of the present invention is as follows.

1. By weight%, C: 0.01-0.25
%, P: 0.03% or less, S: 0.015% or less, C
r: less than 0.5 to 8%, V: 0.05 to 0.5%, N
b: 0 to 0.2%, Ti: 0 to 0.1%, Ta: 0 to 0%
0.2% and N: 0 to 0.1%, and spherical or disk-shaped matched precipitates having a diameter of 2 to 30 nm are present at a density of 1 / μm ³ or more in the crystal grains, and are cementite. ,
One or more types of grain boundary precipitates of M ₇ C ₃ type carbide and M ₂₃ C ₆ type carbide are present, each of which contains 2% or more of V by weight and Ferritic heat-resistant steel with excellent high-temperature strength having a ratio of major axis to minor axis (minor axis / major axis) of 0.5 or more.

2. After normalizing the ferritic heat-resistant steel having the chemical composition described in the above (1) in a temperature range of (T-200) to (T + 200) ° C, the steel is cooled at a speed of 200 ° C / hour or more, and then A method for producing a ferritic heat-resistant steel excellent in high-temperature strength, which is subjected to a tempering treatment in a temperature range of 550 to 850 ° C.

Here, T is a value satisfying the following equation (1).

T = 10 ⁴ /{1.52(A−B)}−273 (1) Here, A and B are values obtained by the following equations (2) and (3). I do. The symbol of the element in the formula represents the content (% by weight) of each element in the steel.

A = 7.0 × ｛V / (V + Nb + Ti + Ta)｝ + 5.2 × ｛(Ti / (V + Nb + Ti + Ta)｝ + 4.8 × ｛(Ta + Nb) / (V + Nb + Ti + Ta)｝ (2) B = log (V + Nb + Ti + Ta) ² (3) The heat-resistant steel of the present invention is mainly a forged steel. And cast steel.

In order to improve the high-temperature strength of low and medium Cr ferritic steels, especially the creep strength at 400 ° C. or higher, the present inventors have focused on precipitation strengthening with precipitates that are stable even at high temperatures. As a result of repeating various tests on the precipitation behavior of carbides and the long-term stability of the underlying structure, the following findings were obtained and the present invention was completed.

(A) Precipitates in low and medium Cr ferritic heat-resistant steels are fine carbonitrides mainly composed of V, Nb, Ti and the like precipitated in crystal grains (hereinafter, the metal element is expressed as M). MX) and coarse carbides precipitated at grain boundaries and lath interfaces.

(B) MX has a spherical shape in the early stage of precipitation at a high temperature, has the same body-centered cubic structure (bcc) as the parent phase, and is in perfect matching with the parent phase. In addition, due to tempering and high-temperature aging that occurs during use, MX has a face-centered cubic structure (fc).
The shape changes to c), and the shape changes to a thin disk shape. However, while the shape is a disk shape, the matching relationship with the parent phase is maintained.

(C) While MX maintains the consistency with the parent phase, dislocations are fixed to the matching strain generated around MX and the dislocations are difficult to move, so that the recovery softening of martensite and bainite is suppressed. It is suppressed and the deformation resistance is increased. Further, dislocations that move during plastic deformation are also fixed, so that deformation resistance increases. Therefore, room temperature strength and creep strength increase.

(D) While the MX maintains consistency with the mother phase, the MX is restrained by the mother phase, so that the growth and aggregation coarsening of the MX itself are suppressed. Therefore, the fine MX is kept stable and at a high density until after the use for a long time at a high temperature, and the precipitation strengthening ability is maintained, so that a stable creep strength can be obtained. That is, the precipitation strengthening ability can be used to the maximum.

(E) As the high-temperature aging proceeds, the shape of MX changes to a cube. When the shape of MX becomes a cube, the matching relationship with the parent phase is completely broken, and the matching strain disappears.

(F) When the consistency between MX and the parent phase is lost, the room temperature strength and the high temperature creep strength decrease.

(G) If the consistency between MX and the matrix is lost, the coarsening of MX proceeds, and the martensite or bainite structure changes significantly with the passage of time. For this reason, stable creep strength cannot be obtained.

(H) However, when the shape, size, and particle density of MX are regulated as follows, the development of high-temperature aging can be suppressed, and the creep strength due to intragranular strengthening can be maintained for a long time.

Shape: spherical or disk-shaped Size: 30 nm or less Particle density: 1 particle / μm ³ or more (i) However, even if the intragranular strengthening is achieved by fine MX, if the grain boundary strength is weak, the creep strength ,
Not only creep ductility but also toughness is deteriorated. Furthermore, tempering embrittlement is likely to occur.

(J) Coarse carbides precipitated at grain boundaries and lath interfaces are M ₂₃ C ₆ , M ₇ C ₃ , cementite, etc., and these relatively coarse carbides precipitate along the grain boundaries, and Precipitates in the form of a film along the grain boundaries.

(K) Around the film-shaped grain boundary precipitate, a non-precipitated zone of other carbides such as MX is formed, the grain boundary strength is weakened, and this appears as a decrease in creep ductility and a deterioration in toughness.

(L) When tempered at a temperature of 550 ° C. or higher, the grain boundary precipitate changes to a spherical shape, and around the spherical grain boundary precipitate, a carbide-free precipitation zone recovers, and accordingly, creep ductility and toughness are increased. Recovers.

(M) When V forms a solid solution in M ₂₃ C ₆ , M ₇ C ₃ , or cementite, coarsening of these carbides is suppressed, and a decrease in creep strength on a long time side is suppressed.

(N) When the ratio of the minor axis to the major axis of M ₂₃ C ₆ , M ₇ C ₃ , or cementite is 0.5 or more, the creep ductility and impact properties are good.

[0041]

Next, the reasons for limiting the chemical composition and precipitates of the heat-resistant steel of the present invention will be described in detail below. In the following description, all percentages of the content of chemical components mean weight%.

C: 0.01 to 0.25% C forms carbides (MX and grain boundary carbides) with Cr, Fe, V, etc., and contributes to high-temperature strength, and itself as an austenite stabilizing element. Stabilize the organization.
The steel of the present invention has a ferrite, martensite, bainite, pearlite, or a mixed structure thereof, and the C content is also important for controlling the balance of these structures. If the C content is less than 0.01%, the amount of carbide precipitation is insufficient, and the hardenability is reduced, resulting in impaired strength and toughness. On the other hand, if it exceeds 0.25%, carbides are excessively precipitated, and the steel is hardened remarkably, impairing workability and weldability. Therefore, the range of the C content is 0.01 to 0.25.
%.

Cr: 0.5 to less than 8% Cr is an essential element for improving oxidation resistance and high-temperature corrosion resistance. If the Cr content is less than 0.5%, these effects cannot be obtained. On the other hand, if it is 8% or more, the weldability and the thermal conductivity are reduced, and the economic efficiency is reduced, and the advantages of the low and medium Cr ferrite steel are reduced. Therefore, the range of the Cr content is set to 0.5% or more and less than 8%.

V: 0.05-0.5% V is an important element forming an MX-type precipitate. That is, V combines with C to form fine VC, and VC is coherently precipitated with the parent phase. However, if it is less than 0.05%, the precipitation amount of VC is small, and does not contribute to the improvement of the room temperature strength and the creep strength. On the other hand, if it is contained in excess of 0.5%, VC does not precipitate in a coherent manner and tends to become coarse, so that the strength and toughness are impaired. Therefore, the V content is set to 0.05 to 0.5%.

P: 0.03% or less, S: 0.015% or less P and S are unavoidable impurity elements, all of which are harmful to toughness, workability and weldability, and particularly promote tempering embrittlement. . For this reason, it is desirable to make it as low as possible, the allowable upper limit of P is 0.03%, and the allowable upper limit of S is 0.0%.
15%.

The heat-resistant steel of the present invention may optionally contain the following alloy elements in addition to the above-mentioned alloy components.

Nb: 0 to 0.2% Nb combines with C similarly to V to form MX [(V, Nb) C] and contributes to improvement in creep strength. Nb has an effect of suppressing the coarsening of MX and increasing the thermal stability of MX, and thus prevents a decrease in creep strength on the long-time side. Further, it is effective for making the crystal grains fine and improving the weldability and toughness. When Nb is contained, the content is preferably 0.002% or more in order to obtain the above effect. On the other hand, 0.
If it exceeds 2%, the steel is hardened remarkably, impairing toughness, workability and weldability. Therefore, when Nb is contained, the content is preferably set to 0.002 to 0.2%.

Ti: 0 to 0.1% Ti, like V, combines with C to form MX and contributes to improvement in creep strength. Furthermore, the crystal grains are refined,
It is also effective in preventing the welding heat affected zone (HAZ) from softening. In order to obtain these effects, the content is preferably 0.001% or more. On the other hand, if it exceeds 0.1%, the steel is hardened remarkably,
Impairs toughness, workability and weldability. Therefore, when Ti is contained, the content is preferably 0.001 to 0.1%.

Ta: 0 to 0.2% Ta, like V, combines with C to form MX and contributes to improvement in creep strength. Furthermore, the crystal grains are refined,
It is also effective in preventing HAZ softening. In order to obtain these effects, the content is preferably 0.002% or more.
If it exceeds 0.2%, the steel is hardened remarkably, impairing toughness, workability and weldability. Therefore, when Ta is contained, the content is preferably 0.002 to 0.2%.

N: 0 to 0.1% N, when combined with V, Nb, Ti, and Ta, forms a fine nitride or a carbonitride compounded with a carbide to improve the creep strength and increase the crystal grain size. Contributes to improvement in toughness due to miniaturization of, and prevention of softening of HAZ. In order to obtain these effects, the content is preferably 0.001% or more. on the other hand,
If the content exceeds 0.1%, the nitride coarsens, impairing strength, toughness, weldability, and workability. In addition, excess N destabilizes the bainite, martensite, and pearlite structures at elevated temperatures. Therefore, in the case where N is contained, 0.
It is good to be 001-0.1%.

Mo: Mo has the effect of solid solution strengthening.
In addition, since it precipitates in combination with VC or the like, it also has the effect of strengthening precipitation and is an element effective for improving creep strength. In order to obtain these effects, the content is preferably 0.01% or more. On the other hand, if the content exceeds 2.5%, the effect is saturated and the weldability and the toughness are impaired. Therefore, M
When making o content, the range is 0.01-2.5.
%.

W: W has a solid solution strengthening effect similarly to Mo. In addition, it is an element effective for improving creep strength because it also has a precipitation strengthening effect due to complex precipitation with VC and the like. In order to obtain these effects, the content is preferably 0.02% or more. . On the other hand, if the content exceeds 5%, the effect is saturated and the weldability and toughness are impaired. Therefore,
When W is contained, the range is preferably set to 0.02 to 5%.

B: B is an element effective for securing stable strength by improving hardenability. To achieve this effect, use 0.
It is preferable to contain 0001% or more. On the other hand, 0.1
%, B segregates remarkably at the grain boundaries, coarsens carbides and causes a decrease in strength and toughness. Therefore, B
When containing, the range is 0.0001 to 0.1
%.

Co: Co is an austenite stabilizing element and has a solid solution strengthening action. In order to obtain these effects, the content is preferably 0.01% or more. on the other hand,
If it exceeds 0.5%, the high-temperature creep strength decreases. In addition, excessive addition is not preferable from the viewpoint of economy. Therefore, when Co is contained, the upper limit is preferably set to 0.5%.

Ni: Ni is an austenite stabilizing element and contributes to improvement in toughness. In order to obtain these effects, the content is preferably 0.01% or more. on the other hand,
If it exceeds 0.5%, the high-temperature creep strength decreases. Also, from the viewpoint of economy, addition of a large amount is not preferable. Therefore, when Ni is contained, the upper limit is preferably set to 0.5%.

Cu: Cu is an austenite stabilizing element and contributes to improvement in thermal conductivity. In order to obtain these effects, the content is preferably 0.01% or more. on the other hand,
If it exceeds 0.5%, the high temperature creep strength and the toughness are deteriorated. Therefore, when Cu is contained, the upper limit is set to 0.1.
It is good to make it 5%.

Al: Al is an element effective as a deoxidizing agent. In order to obtain this effect, the content is preferably 0.001% or more. On the other hand, when it exceeds 0.05%, creep strength and workability are impaired. Therefore, when Al is contained, the range is preferably made 0.001 to 0.05%. A desirable upper limit is 0.02%.

Si: Si is an element that acts as a deoxidizing agent and enhances the steam oxidation resistance of steel. In order to obtain these effects, the content is preferably 0.01% or more. On the other hand, the toughness exceeding 0.7% is significantly reduced, and is harmful to the creep strength. Therefore, when Si is contained, the range is preferably set to 0.01 to 0.7%. A desirable upper limit is 0.5%.

Mn: Mn improves hot workability by the desulfurization and deoxidation effects during melting. Furthermore, hardenability is improved. In order to obtain these effects, the content is preferably 0.01% or more. On the other hand, if it exceeds 1%, the stability of the fine carbonitride effective for creep strengthening is impaired, and the creep strength at high temperature and long time decreases. Therefore, when Mn is contained, the content is preferably 0.01 to 1%. A desirable upper limit is 0.5%.

Next, the precipitate will be described.

Coordinated precipitates: Coordinated precipitates are carbonitrides that are deposited in conformity with the parent phase, and a coherent strain occurs around the precipitates. Therefore, whether the precipitate is a consistent precipitate can be determined by examining the presence or absence of a conformational strain by observation with a transmission electron microscope. Specifically, by performing a two-wave approximation method at a high magnification of 20,000 or more using a transmission electron microscope, it is possible to determine the presence or absence of a matching distortion that appears as a matching distortion contrast.
The matched precipitate is an MX-type precipitate, where MX is VC, V
N, NbC, NbN, TiC, TiN, TaC, TaN
Etc. and these composite precipitates.

[M is a metal atom (mainly V, Nb, T
i, Ta), X is C or N] These are coherently precipitated in the parent phase and increase the room-temperature strength and creep strength due to coherent strain occurring around them. When they are spherical or disk-shaped, they are coherently precipitated with the parent phase, and thus have the effect of increasing the strength. However, when they are cubic, the coherence relationship with the parent phase is broken, and the strength effect disappears.

If the diameter of the coherent precipitate exceeds 30 nm in diameter, the dislocation fixing force is reduced and the strengthening effect is reduced. Therefore, the upper limit is set to 30 nm. The resolution of a normal transmission electron microscope (eg, acceleration voltage 200 kv) is 2n
The size of MX to be counted is set to 2 nm or more, since the size is smaller than m and it is difficult to confirm the size smaller than m. Further, even if a consistent precipitate having a diameter of 30 nm or less is precipitated, if the particle density is less than 1 / μm ^3, the strengthening ability is not sufficient, so that the particle density is 1 / μm ³ or more.

The particle density is determined, for example, by using a transmission electron microscope under two-wave approximation diffraction conditions, then observing at 50000 fields or more at a magnification of 40,000 or more, and measuring the number of particles having matching contrast.

Grain boundary precipitate: M ₂₃ C ₆ , M ₇
C ₃ or cementite (Fe ₃ C), and the heat-resistant steel of the present invention contains one or more of these. The grain boundary strength is improved by setting the ratio of the minor axis to the major axis (minor axis / major axis) of these grain boundary precipitates to 0.5 or more. If this ratio is less than 0.5, the grain boundary strength will be reduced, and the toughness and creep ductility will be significantly reduced. Therefore, the ratio of the minor axis to the major axis of the grain boundary precipitate is set to 0.5 or more.

When the V content of the metal atoms constituting M ₂₃ C ₆ , M ₇ C ₃ or cementite is less than 2% by weight, the coarsening of these carbides is promoted and the creep strength is reduced. descend. Therefore, the V content in M ₂₃ C ₆ , M ₇ C ₃ or cementite is 2% or more by weight%. The upper limit is not particularly limited. However, if V is precipitated as M ₂₃ C ₆ , M ₇ C ₃ or cementite, the amount of MX deposited decreases, so that it is preferably 10% or less. In addition, M ₂₃ C _6, M ₇
The amount of V in C ₃ or cementite is determined by EDX (energy dispersive X-ray spectroscopy) of a transmission electron microscope.

Hereinafter, the manufacturing method will be described.

The heat-resistant steel to be subjected to the production method of the present invention is:
Forged steel that has been melted, cast and hot worked, and cast steel used as cast. By subjecting the steel having the chemical composition specified in the present invention to the following heat treatment, the matched precipitates and the grain boundary precipitates can be set to predetermined sizes, shapes, compositions, and densities.

(1) Normalization Normalization is preferably performed at a temperature equal to or higher than the austenite transformation start temperature and at a temperature at which MX forms a solid solution. The undissolved MX is coarse and is a non-coherent precipitate, and therefore has no creep strengthening effect. When the undissolved MX increases, the creep strength and toughness are reduced. In addition, undissolved M
As the amount of X increases, the precipitation density of finely matched MX precipitated during subsequent tempering or long-term aging decreases, and a sufficient strengthening action cannot be obtained. Specifically, in normalizing at a temperature of (T-200) ° C. or lower, the amount of undissolved MX increases and the toughness deteriorates.

Here, T is a temperature obtained by the following equation (1). Equation (1) is an empirical equation obtained by various experiments.

T = 10 ⁴ /{1.52(A−B)}−273 (1) Here, A and B are values obtained by the following equations (2) and (3). I do. A = 7.0 × ｛V / (V + Nb + Ti + Ta)｝ + 5.2 × ｛(Ti / (V + Nb + Ti + Ta)｝ + 4.8 × ｛(Ta + Nb) / (V + Nb + Ti + Ta)｝ ・ ・ ・ ・ ・ ・ (2) B = logｌｏ (V + Nb + Ti + Ta) ² ｝ ・ ・ ・ (3) Also, if normalizing treatment is performed at a temperature exceeding (T + 200) ° C. However, the precipitation form and composition of carbide satisfy the conditions specified in the present invention, but the crystal grains become coarse and the toughness is deteriorated, so that the normalizing temperature range is (T-200) to (T + 2).
00) ° C.

(2) Cooling after normalizing If cooling is performed at a low speed of 200 ° C./hour or less, coarse MX and other carbides precipitate during cooling, and sufficient hardenability cannot be obtained. During cooling, P and S segregate at the grain boundaries, and tempering embrittlement and creep embrittlement are likely to occur. Therefore, the cooling rate after normalizing was set to 200 ° C./hr or more. The higher the cooling rate, the better. However, practically, the cooling rate is preferably 5 ° C./sec or less, which is a cooling rate corresponding to water cooling.

(3) Tempering Tempering is essential for precipitating MX carbide.
However, when the heating temperature is 550 ° C. or lower, MX does not sufficiently precipitate, and the strengthening ability is small. Further, M ₂₃ C ₆ , M ₇ C ₃ , and cementite precipitate in the form of a film at the grain boundaries, thereby deteriorating the toughness. On the other hand, tempering at 850 ° C. or more causes the MX to become coarse and deteriorates the strength and toughness. Therefore, the tempering temperature was 550 ° C to 850 ° C.

As described above, the manufacturing method of the present invention can be applied to forged steel and cast steel. However, in a forged steel hot-rolled in a high-temperature austenite region, many dislocations are introduced. The particle density of the matched MX increases, and it is easy to increase the strength. This is because dislocations serve as nucleation sites for precipitation, and the precipitation density increases when many dislocations are introduced.

However, in order to take advantage of the effect of hot working, A
It is preferable to heat to a temperature between c3 point and 1300 ° C. and perform hot working at a reduction of 50% or more. If any one of the heating temperature and the rolling reduction is not in the above range, no dislocation is accumulated, and the effect of hot working does not appear. Further, if the normalizing is performed continuously without lowering the temperature below the normalizing temperature after the hot working, the manufacturing cost can be reduced by energy saving.

[0076]

EXAMPLE In a 150 kg vacuum melting furnace, 19 types of steels having the chemical compositions shown in Table 1 were melted and cast to obtain an ingot.
Forged in the temperature range of 200 to 1000 ° C and thickness 50mm
Plate.

[0077]

[Table 1]

Regarding steel symbol 1, a steel plate was manufactured by directly machining from an ingot without forging.

[0079]

[Table 2]

[0080]

[Table 3]

From the tempered steel sheet, a thin film sample was prepared by electropolishing, and the thin film sample was prepared using a transmission electron microscope (acceleration voltage 200 k
Observed by V), the size, particle density and shape of the precipitate were measured.

Here, the matched precipitates were determined by the two-wave approximation observation method using a transmission electron microscope based on the presence or absence of matched strain contrast. The average diameter and particle density of the matched precipitates were measured by injecting an electron beam perpendicular to the parent phase {001}. The flattening rate of the grain boundary precipitate was arranged by the ratio of the minor axis to the major axis. The amount of V in the grain boundary precipitate was measured by EDX analysis of the precipitate observed with a transmission electron microscope.

In the creep test, a test piece having a diameter of 6 mm and a length of a parallel portion of 30 mm was manufactured, and the test pieces were formed at 500 ° C. and 550 ° C.
Test up to 10,000 hours at
The average creep rupture strength at 0 ° C. × 8000 hours was determined.

The rate of decrease in strength due to long-term creep was quantified by organizing the ratio of the breaking strength at 10000 hours to the breaking strength at 10 hours at each temperature.

In the Charpy impact test, 10 × 10 × 55
(Mm) 2 mm V notch test piece (JIS No. 4 test piece)
Was used to determine the ductile-brittle fracture transition temperature.

Tables 4 and 5 show the evaluation results.

[0087]

[Table 4]

[0088]

[Table 5]

The following is clear from Tables 4 and 5.

Even if the chemical composition falls within the range specified by the present invention as in the steel symbols 1 to 5 of the comparative examples, if the conditions deviate from the conditions specified by the heat treatment method, the size of the coherent precipitate,
The particle density and the shape of the grain boundary precipitate did not fall within the specified ranges, and the creep strength or the Charpy impact characteristics were significantly deteriorated.

Further, in the case of steel symbol A having an excessive C content, the shape of the grain boundary precipitate became a film, and the toughness was deteriorated. Since the carbides were coarsened by the long-term aging, the creep strength on the long-term side was significantly reduced.

[0092] In the steel symbol B containing no Cr, the structure was unstable at high temperatures and the strength was significantly reduced.

In steel symbol C having a low V content, no coherent precipitate was precipitated and the creep strength was low.

In the steel D containing excessive V, the coherent precipitates became coarse and the particle density was lowered. Therefore, sufficient creep strength could not be obtained, resulting in a significant decrease in toughness.

In the steel symbol E containing excessive amounts of Nb, N and Ni, coarse nitrides were precipitated and the toughness was deteriorated.

A steel symbol F in which Ti and Mo are excessively added is
In addition to the precipitation of coarse carbonitrides, the carbides tended to become coarse at high temperatures, and the creep strength and toughness were significantly reduced.

In steel symbol G containing excessive Mn, carbides tend to coarsen at a high temperature, and the strength was significantly reduced at a long time of 550 ° C.

In the H steel containing excessive amounts of Ta, W and Co, the strength is unstable at a high temperature and the coarse carbonitride precipitates, so that the creep strength and the toughness are remarkably deteriorated.

[0099] In the case of steel symbol I in which B is excessive, coarse carbides are precipitated in the form of a film at the grain boundaries, so that the creep strength and the toughness are significantly reduced.

On the other hand, in the steel obtained by the production method of the present invention, fine consistent precipitates are deposited at a high density,
The grain boundary precipitate has a flattening rate of 0.65 or more (flattening rate of 1.0
Was spherical). The creep strength at 500 ° C. for 8000 hours was as high as 236 MPa or more.
The ratio of the creep strength at 10000 hours / 10 hours at 00 ° C. and 550 ° C. was as high as 0.57 or more and 0.47 or more, respectively. This indicates that the strength is not significantly reduced by long-time creep. Further, the ductile-brittle wavefront transition temperature was -27 ° C or less, and all showed good toughness.

[0101]

According to the present invention, a low and medium Cr ferritic heat-resistant steel having a stable and high creep rupture strength and excellent toughness even when used at a high temperature of 400 ° C. or more for a long time can be obtained. It can also be used in applications where high Cr ferritic steels have been used in the past, with great economical effects.

Claims (2)

[Claims] (1) C: 0.01 to 0.25% by weight%
P: 0.03% or less, S: 0.015% or less, Cr:
0.5 to less than 8%, V: 0.05 to 0.5%, Nb: 0
-0.2%, Ti: 0-0.1%, Ta: 0-0.2%
And N: 0 to 0.1%, and the crystal grains have a diameter of 2%.
1 spherical / disk-shaped matched precipitate of ~ 30 nm / μ
present in m ³ or more density and there cementite, one or more grain boundary precipitates of M ₇ C ₃ -type carbide and M ₂₃ C ₆ type carbide, weight both their grain boundary precipitates A ferritic heat-resistant steel excellent in high-temperature strength, characterized by containing V of 2% or more in% and having a ratio of minor axis to major axis (minor axis / major axis) of 0.5 or more. 2. After the ferritic heat-resistant steel having the chemical composition according to claim 1 is tempered in a temperature range of (T-200) to (T + 200) ° C., the rate is increased at a rate of 200 ° C./hour or more. A method for producing a ferritic heat-resistant steel having excellent high-temperature strength, comprising cooling and then performing a tempering treatment in a temperature range of 550 to 850 ° C. Here, T is a value satisfying the following equation (1). T = 10 ⁴ /{1.52(A−B)}−273 (1) Here, A and B are values obtained by the following equations (2) and (3). The symbol of the element in the formula represents the content (% by weight) of each element in the steel. A = 7.0 × ｛V / (V + Nb + Ti + Ta)｝ + 5.2 × ｛(Ti / (V + Nb + Ti + Ta)｝ + 4.8 × ｛(Ta + Nb) / (V + Nb + Ti + Ta)｝ ・ ・ ・ ・ ・ ・ (2) B = log (V + Nb + Ti + Ta) ²・ ・ ・ ・ (3)