CN112585288A

CN112585288A - Austenitic steel sintered material, austenitic steel powder and turbine component

Info

Publication number: CN112585288A
Application number: CN201980050604.8A
Authority: CN
Inventors: 芝山隆史; 今野晋也
Original assignee: Mitsubishi Power Ltd
Current assignee: Mitsubishi Power Ltd
Priority date: 2018-11-19
Filing date: 2019-11-11
Publication date: 2021-03-30
Also published as: KR102467393B1; SG11202100355UA; EP3653322A1; JP2020084229A; WO2020105496A1; KR20210024083A; US20200157664A1; JP7112317B2

Abstract

The purpose of the present invention is to provide an austenitic steel sintered material, an austenitic steel powder, and a turbine component that have strength equal to or higher than that of a Ni-based alloy and are less susceptible to oxygen. The present invention provides an austenitic steel sintered material and an austenitic steel powder, which contain, in mass%, Ni: 25-50%, Cr: 12-25%, Nb: 3-6%, B: 0.001 to 0.05%, Ti: 1.6% or less, W: 6% or less, Mo: 4.8% or less, Zr: 0.5% or less, and the balance of Fe and inevitable impurities.

Description

Austenitic steel sintered material, austenitic steel powder and turbine component

Technical Field

The invention relates to an austenitic steel sintered material, an austenitic steel powder and a turbine component.

Background

Recently, the steam temperature has been raised to a high temperature with the aim of increasing the efficiency of coal-fired power plants. In a currently operated coal-fired power plant, a steam turbine (USC: Ultra Super Critical pressure power generation) having the highest steam temperature is operated at a steam temperature of 620 ℃ stage, but the CO is suppressed₂The discharge is further elevated in consideration of the future. Heretofore, as high-temperature components of steam turbines, 9 Cr-based and 12 Cr-based heat-resistant ferritic steels and the like have been used, but it is considered that the application of these steels becomes difficult as the steam temperature increases.

As an alloy suitable for high-temperature members, Ni-based alloys having a higher durability temperature than ferritic steels may be a candidate. The Ni-based alloy uses Al and Ti as precipitation strengthening elements, generates a γ' phase as a stable phase at high temperatures, and exhibits excellent strength at high temperatures. However, although the turbine valve housing, the turbine disk, and the like are generally manufactured by a casting method, the casting method is insufficient in blocking from air during melting, and if the active elements (Al and Ti) are large, these elements are oxidized.

Patent document 1 discloses a technique of applying, to a turbine component, an austenitic steel having both excellent strength and castability in place of a Ni-based alloy, and an austenitic steel casting product using the austenitic steel.

Documents of the prior art

Patent document

Patent document 1: japanese laid-open patent publication No. 2017-88963

Disclosure of Invention

Problems to be solved by the invention

The above patent document 1 proposes a composition of austenitic steel in which macroscopic defects are reduced in a large-sized cast product, but the production of a mold used for the cast product is relatively troublesome. In particular, if the mold is a large mold for a casting product having a complicated shape, the process cost increases. Therefore, if a member can be obtained by sintering without casting, the manufacturability of the turbine member can be further improved.

In view of the above circumstances, an object of the present invention is to provide an austenitic steel sintered material, an austenitic steel powder, and a turbine component, which have strength equal to or higher than that of a Ni-based alloy and are less susceptible to oxygen.

Means for solving the problems

A first aspect of the present invention for solving the above problems is an austenitic steel sintered material containing, in mass%, Ni: 25-50%, Cr: 12-25%, Nb: 3-6%, B: 0.001 to 0.05%, Ti: 1.6% or less, W: 6% or less, Mo: 4.8% or less, Zr: 0.5% or less, and the balance of Fe and inevitable impurities.

A second aspect for solving the above problems is an austenitic steel sintered material containing, in mass%, Ni: 30-45%, Cr: 12-20%, Nb: 3-5%, B: 0.001 to 0.02%, Ti: 0.3-1.3%, W: 5.5% or less, Mo: 2% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

A third aspect for solving the above problems is an austenitic steel sintered material containing, in mass%, Ni: 30-40%, Cr: 15-20%, Nb: 3.5-4.5%, B: 0.001 to 0.02%, Ti: 0.5-1.1% or less, W: 5.5% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

A fourth aspect for solving the above problems is a turbine component using an austenitic steel sintered material.

A fifth aspect of the present invention for solving the above problems is an austenitic steel powder containing, in mass%, Ni: 25-50%, Cr: 12-25%, Nb: 3-6%, B: 0.001 to 0.05%, Ti: 1.6% or less, W: 6% or less, Mo: 4.8% or less, Zr: 0.5% or less, and the balance of Fe and inevitable impurities.

A sixth aspect for solving the above problems is an austenitic steel powder containing, in mass%, Ni: 30-45%, Cr: 12-20%, Nb: 3-5%, B: 0.001 to 0.02%, Ti: 0.3-1.3%, W: 5.5% or less, Mo: 2% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

A seventh aspect for solving the above problems is an austenitic steel powder containing, in mass%, Ni: 30-40%, Cr: 15-20%, Nb: 3.5-4.5%, B: 0.001 to 0.02%, Ti: 0.5-1.1% or less, W: 5.5% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

More specific configurations of the present invention are described in the claims.

Effects of the invention

According to the present invention, it is possible to provide an austenitic steel sintered material, an austenitic steel powder, and a turbine component that have strength equal to or higher than that of a Ni-based alloy and are less susceptible to oxygen.

Problems, configurations, and effects other than those described above will become apparent from the following description of the embodiments.

Drawings

FIG. 1A is a schematic view showing an example of the structure of a sintered austenitic steel material obtained by sintering an austenitic steel powder according to the present invention.

FIG. 1B is an SEM observation photograph showing an example of the structure of a sintered austenitic steel material obtained by sintering the austenitic steel powder of the present invention.

Fig. 2 is a schematic view of an example of the structure of the austenitic steel casting material of patent document 1.

FIG. 3 is a schematic view showing an example of the structure of a conventional Ni-based alloy forging material.

FIG. 4 is a schematic view showing an example of a turbine valve housing to which an austenitic steel sintered material obtained by sintering an austenitic steel powder according to the present invention is applied.

FIG. 5 is a schematic view showing an example of a turbine disk to which an austenitic steel sintered material obtained by sintering an austenitic steel powder according to the present invention is applied.

FIG. 6 is a graph showing the 0.2% yield strength ratio (based on comparative example 4) of examples 1 to 3 and comparative examples 1 to 4.

FIG. 7 is a graph showing creep rupture temperature ratios (based on comparative example 3) of examples 1 to 3 and comparative examples 1 to 4.

FIG. 8 is a graph showing the 0.2% yield strength ratio and creep rupture temperature ratio of examples 1 and 3 and comparative examples 1, 3 and 4.

Detailed Description

Hereinafter, the present invention will be described in detail with reference to the accompanying drawings.

[ Austenitic steel powder and Austenitic steel sintered Material ]

FIG. 1A is a schematic view showing an example of the structure of a sintered austenitic steel material obtained by sintering an austenitic steel powder according to the present invention, and FIG. 1B is an SEM observation photograph showing an example of the structure of a sintered austenitic steel material obtained by sintering an austenitic steel powder according to the present invention. As shown in fig. 1A and 1B, the austenitic steel sintered material of the present invention has austenitic steel powder crystals 1, grain boundaries 2 existing at the boundaries of adjacent austenitic steel crystals, and a laves (laves) phase 3 precipitated on the grain boundaries 2.

The average grain size of the austenitic steel powder crystals 1 is preferably 10 to 300. mu.m. If the thickness is less than 10 μm, the creep strength may become insufficient. If it exceeds 300. mu.m, the tensile strength and fatigue strength may become insufficient. In addition, since the total amount of grain boundaries changes, the grain boundary coverage of the lafves phase changes, and the strength (creep strength, tensile strength, fatigue strength, and the like) may decrease. The "particle diameter" can be measured as a planar image when observed by an observation means such as an electron microscope. The "average grain size" may be a value obtained by averaging the grain sizes of a predetermined number of austenite steel powder crystals 1 shown in an observation photograph of a predetermined magnification.

As a comparison with the above-described structure of the austenitic steel sintered material of the present invention, the cast structure of patent document 1 and the structure of a conventional Ni-based Alloy (Alloy718) will also be described. Fig. 2 is a schematic view of an example of the structure of the austenitic steel casting material of patent document 1. As shown in fig. 2, the austenitic steel casting material has austenitic steel powder crystals 4, grain boundaries 5 existing at the boundaries of adjacent austenitic steel crystals, and a lafves phase 6 precipitated on the grain boundaries 5. The cast structure has few grain boundaries, and the crystal has uneven grain size and shape. Further, the cast structure has a larger micro-segregation than the structure of the sintered material. It is considered that the larger the member is, the larger the micro segregation is, and defects due to the micro segregation are likely to be generated, thereby reducing the strength. On the other hand, since the sintered material forms a homogeneous structure regardless of the size of the member, micro-segregation is less likely to occur.

FIG. 3 is a schematic view showing an example of the structure of a conventional Ni-based Alloy forging material (Alloy 718). As shown in fig. 3, the Ni-based alloy forged material has Ni-based alloy crystals 7, old grain boundaries (PPB)8 existing at the boundaries of adjacent Ni-based alloy crystals, and a δ phase 9 precipitated on the old grain boundaries (PPB) 8.

As is clear from comparison of FIGS. 1 to 3, the structure of the sintered austenitic steel of the present invention is clearly distinguished from the structure of the conventional cast austenitic steel and Ni-based alloy.

The composition of the austenitic steel sintered material of the present invention will be described below. In the following description of the composition, "%" means "% by mass" unless otherwise specified.

Ni (nickel): 25 to 50 percent

Ni is added as an austenite phase stabilizing element. Further, an intermetallic compound (delta phase, Ni) is formed with Nb described later₃Nb) and precipitates in the grains, contributing to the intragranular strengthening. From the viewpoint of phase stability, Ni is preferably 25 to 50% (25% to 5%), more preferably 30 to 45%, and still more preferably 30 to 40%.

Cr (chromium): 12 to 25 percent

Cr is an element that improves oxidation resistance and steam oxidation resistance. Sufficient oxidation resistance can be obtained by adding 12% or more in consideration of the operating temperature of the steam turbine. When more than 25% is added, intermetallic compounds such as σ phase precipitate, resulting in a decrease in high-temperature ductility and toughness. In view of the balance thereof, the amount of Cr is preferably 12 to 25%, more preferably 12 to 20%, and still more preferably 15 to 20%.

Nb (niobium): 3 to 6 percent

Nb for Laves phase (Fe)₂Nb) and delta phase (Ni)₃Nb) is added for stabilization. As shown in fig. 2, the lafves phase 6 is mainly precipitated at the grain boundaries 2, and contributes to grain boundary strengthening. The delta phase is mainly precipitated in the grains and contributes to strengthening. By adding 3% or more, sufficient high-temperature creep strength can be obtained. If more than 6% is added, the δ phase may be adversely compatible and easily separated out. In order to obtain a sufficient strength, the amount of Nb is preferably 3 to 6%, more preferably 3 to 5%, and still more preferably 3.5 to 4.5%.

B (boron): 0.001 to 0.05 percent

B contributes to the precipitation of the lafves phase at the grain boundaries. When B is not added, the Laves phase in the grain boundary is hard to precipitate, and the creep strength and creep ductility are reduced. By adding 0.001% or more, the effect of grain boundary precipitation can be obtained. On the other hand, if the amount is too large, the melting point is locally lowered, and there is a possibility that, for example, weldability is lowered. In view of this, the amount of B is preferably 0.001 to 0.05%, more preferably 0.001 to 0.02%.

Ti (titanium): 0 to 1.6 percent

Ti is an element contributing to precipitation strengthening in grains of the γ "phase and the δ phase. By adding the additive appropriately, the initial creep deformation can be greatly reduced. However, if the amount is excessively added, the mechanical properties are adversely affected by oxidation during production. In view of this, Ti is preferably 1.6% or less, more preferably 0.3 to 1.3%, and further preferably 0.5 to 1.1%.

W (tungsten): 0 to 6 percent

W contributes to stabilization of the lafuties phase in addition to solid solution strengthening. The addition of W increases the amount of laves phase precipitated at grain boundaries, and contributes to the improvement of fracture strength and ductility in creep characteristics over a long period of time. If the amount exceeds 6%, harmful compatibility such as the δ phase may be easily precipitated. In view of this, W is preferably 6% or less, more preferably 5.3 to 6% or less, and further preferably 5.5 to 5.5%.

Mo (molybdenum): 0 to 4.8 percent

Mo contributes to stabilization of the laffse phase in addition to solid solution strengthening. By adding Mo, the amount of precipitates of the lafutiform phase that precipitate at grain boundaries increases, and the lafutiform phase can contribute to fracture strength and ductility in creep characteristics over a long period of time. In view of this, Mo is preferably 0 to 4.8%, more preferably 0 to 2% or less.

Zr (zirconium): 0 to 0.5 percent

As with B, the addition of Zr contributes to the precipitation of a Laves phase at grain boundaries and also contributes to a gamma' phase (Ni)₃Nb) is precipitated. It is particularly effective in a short time or at a low temperature (less than 750 ℃ C., preferably 700 ℃ C. or less). However, since the γ "phase is a quasi-stable phase, it changes to the δ phase when it is held at a high temperature (particularly 750 ℃ or higher) for a long time. Therefore, it may not be added. If the amount is too large, the stability of the delta phase is improved, and the gamma phase is rapidly changed to the delta phase. In addition, weldability deteriorates. In view of this, Zr is preferably 0 to 0.5%, more preferably 0 to 0.3% or less.

The austenitic steel sintered material of the present invention contains Nb and Ti as main strengthening elements, and does not contain Al as a strengthening element, as described above. Therefore, the steel sheet is less susceptible to oxidation or the like by oxygen, and can have improved strength (creep strength, tensile strength, fatigue strength, or the like).

Further, the sintered material has a forged structure, and the strength characteristics can be easily controlled in accordance with the required strength of the product by controlling the crystal grain size by heat treatment or the like.

Further, since the mold of the sintered material is easier to manufacture than the mold of the cast material, it can be manufactured with a high yield even in a complicated product shape.

[ method for producing sintered Austenitic Steel Material ]

Next, a method for producing the sintered austenitic steel material of the present invention will be described. The austenitic steel sintered material of the present invention can be produced, for example, by the following steps.

(1) The raw material powder or the raw material alloy having the above composition is made into an alloy powder (austenitic steel powder) having an average particle diameter of 250 μm or less by a gas atomization method or a water atomization method.

(2) The alloy powder obtained in (1) above was sintered by Hot Isostatic Pressing (HIP). The sintering conditions are, for example, sintering temperature: 1100-1300 ℃ isostatic pressing: above 50 MPa.

Sintering may be performed by hot pressing under a different pressure or by a metal powder injection molding (MIM) method instead of HIP. Further, after sintering, solution heat treatment (heat treatment temperature: 1100 to 1300 ℃) and aging heat treatment (heat treatment temperature: 1000 ℃ or lower) may be performed.

[ turbine component Using sintered Austenitic Steel Material ]

FIG. 4 is a schematic view showing an example of a turbine valve housing to which the sintered austenitic steel material of the present invention is applied, and FIG. 5 is a schematic view showing an example of a turbine disk to which the sintered austenitic steel material of the present invention is applied. As shown in fig. 4, the austenitic steel sintered material of the present invention has excellent strength and is therefore suitable for the turbine valve housing 10 and the turbine disk 11.

Examples

The present invention will be described in more detail below with reference to examples.

[ production and evaluation of sintered Austenitic Steel Material ]

The sintered materials of examples 1 to 3 and comparative examples 1 to 2 were produced and evaluated. The compositions of examples 1 to 3 and comparative examples 1 to 2 are shown in Table 1 below. A master ingot or a raw material having a composition shown in Table 1 was prepared, and alloy powder having a particle size of 250 μm or less was prepared by a gas atomization method. The obtained alloy powder was sintered by HIP (sintering temperature: 1160 ℃ C., isostatic pressure: 100MPa) to prepare sintered materials of examples 1 to 3 and comparative examples 1 to 2. Comparative example 1 had a composition with an amount of Cr outside the range of the present invention, and comparative example 2 had a composition with an amount of Ni outside the range of the present invention.

Alloy (inconel)718 (wrought material) as a Ni-based alloy and alloy (inconel)625 (cast material) as a Ni-based alloy were also prepared as comparative examples 3 and evaluated as comparative examples 4. The compositions of comparative examples 3 and 4 are also shown in Table 1. "INCONEL" is a registered trademark of Huntington Alloys Corporation.

[ Table 1]

	Fe	Ni	Cr	Nb	B	Ti	W	Mo	Zr	C	Al	Manufacturing method
													Example 1	Bal.	35.4	18.3	4.2	0.006	0.76	-	-	0.014	-	-	HIP
Example 2	Bal.	38.0	15.9	4.0	0.008	0.93	-	-	0.016	-	-	HIP
													Example 3	Bal.	36.1	18.3	4.1	0.006	0.84	5.0	-	0.016	-	-	HIP
Comparative example 1	Bal.	35.6	10.8	3.8	0.006	0.78	-	-	0.018	-	-	HIP
													Comparative example 2	Bal.	24.4	19.2	4.3	0.006	0.81	-	-	0.015	-	-	HIP
Comparative example 3	Bal.	52.3	19.0	5.1	0.004	0.97	-	3.1	-	0.03	0.59	Forging
													Comparative example 4	4.0	Bal.	21.8	3.5	-	0.22	-	8.9	-	0.03	0.26	Forging

For examples 1 to 3 and comparative examples 1 to 4, 0.2% yield strength and creep endurance temperature were evaluated. The 0.2% yield strength was tested in accordance with JIS G0567, and the creep test was tested in accordance with JIS Z22761.

FIG. 6 is a graph showing the 0.2% yield strength ratio (based on comparative example 4) of examples 1 to 3 and comparative examples 1 to 4. As shown in fig. 6, the sintered materials of examples 1 and 3 each showed a value higher than that of comparative examples 1, 2 and 4, and showed a 0.2% proof stress ratio equal to or higher than that of conventional comparative example 3(Alloy 718).

FIG. 7 is a graph showing creep rupture temperature ratios (based on comparative example 3) of examples 1 to 3 and comparative examples 1 to 4. As shown in FIG. 7, the sintered materials of examples 1 to 2 all showed higher values than those of comparative examples 1 to 3, and showed 0.2% proof stress ratios equal to or higher than that of conventional comparative example 4(Alloy 625).

From fig. 6 and 7, the 0.2% yield strength ratio of example 2 is slightly lower than that of comparative examples 2 to 4, but the creep rupture temperature is higher than that of comparative examples 2 to 4, and when both the 0.2% yield strength ratio and the creep rupture temperature are combined and judged, it can be said that the creep rupture temperature is superior to that of comparative examples.

Further, from fig. 6 and 7, the creep rupture temperature of example 3 is slightly lower than that of comparative example 4, but the 0.2% yield strength ratio is much higher than that of comparative example 4, and when both the 0.2% yield strength ratio and the creep rupture temperature are combined and judged, it can be said that the creep rupture temperature is superior to that of the comparative example.

FIG. 8 is a graph showing the 0.2% yield strength ratio and creep rupture temperature ratio of example 3 and comparative examples 1, 3 and 4. As shown in FIG. 8, the 0.2% yield strength ratio and the creep rupture temperature ratio of examples 1 and 3 both showed larger values than those of comparative example 1. In addition, with respect to the 0.2% yield strength ratio, examples 1 and 3 were larger than comparative example 4(Alloy625), and achieved the same level as comparative example 3(Alloy 718). Further, examples 1 and 3 are larger than comparative example 3(Alloy718) with respect to the creep endurance temperature ratio. In particular, example 1 achieved a level equivalent to that of comparative example 4(Alloy 625).

The 0.2% yield strength and creep resistance temperature generally show a trade-off relationship, that is, a behavior that the creep resistance temperature becomes lower when the 0.2% yield strength becomes higher, and the 0.2% yield strength becomes lower when the creep resistance temperature becomes higher. Since example 1 and example 3 are located at the upper right position of the straight line connecting comparative example 3 and comparative example 4, if both the 0.2% yield strength ratio and the creep rupture temperature are determined in total, it can be said that they are superior to comparative example 3 and comparative example 4.

As described above, according to the present invention, it has been revealed that an austenitic steel sintered material and a turbine component having strength equal to or higher than that of a Ni-based alloy and less susceptible to oxygen can be provided.

The present invention is not limited to the above embodiments, and includes various modifications. For example, the above-described embodiments have been described in detail to explain the present invention easily and understandably, and are not limited to the embodiments having all the configurations described. In addition, a part of the configuration of one embodiment may be replaced with the configuration of another embodiment, and the configuration of another embodiment may be added to the configuration of one embodiment. Further, a part of the configuration of each embodiment can be added, deleted, or replaced with another configuration.

Description of the symbols

1. 4: austenitic steel powder crystal, 2, 5: grain boundaries, 3, 6: laves phase, 7: ni-based alloy crystal, 8: old particle boundary (PPB), 9: delta phase, 10: turbo valve housing, 11: a turbine disk.

Claims

1. An austenitic steel sintered material, comprising, in mass%, Ni: 25-50%, Cr: 12-25%, Nb: 3-6%, B: 0.001 to 0.05%, Ti: 1.6% or less, W: 6% or less, Mo: 4.8% or less, Zr: 0.5% or less, and the balance of Fe and inevitable impurities.

2. An austenitic steel sintered material, comprising, in mass%, Ni: 30-45%, Cr: 12-20%, Nb: 3-5%, B: 0.001 to 0.02%, Ti: 0.3-1.3%, W: 5.5% or less, Mo: 2% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

3. An austenitic steel sintered material, comprising, in mass%, Ni: 30-40%, Cr: 15-20%, Nb: 3.5-4.5%, B: 0.001 to 0.02%, Ti: 0.5-1.1% or less, W: 5.5% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

4. The sintered austenitic steel material as claimed in any one of claims 1 to 3, wherein the mean grain size of the sintered austenitic steel material is 10 to 300 μm.

5. The sintered austenitic steel material as claimed in any one of claims 1 to 3, wherein a Laves phase is precipitated at grain boundaries of the sintered austenitic steel material.

6. The austenitic steel sintered material of claim 5, wherein the Laves phase is composed of Fe₂And Nb.

7. A turbine component using the austenitic steel sintered material as claimed in any 1 of claims 1 to 3.

8. The turbine component of claim 7, wherein the turbine component is a turbine valve housing or a turbine disk.

9. An austenitic steel powder, comprising, in mass%, Ni: 25-50%, Cr: 12-25%, Nb: 3-6%, B: 0.001 to 0.05%, Ti: 1.6% or less, W: 6% or less, Mo: 4.8% or less, Zr: 0.5% or less, and the balance of Fe and inevitable impurities.

10. An austenitic steel powder, comprising, in mass%, Ni: 30-45%, Cr: 12-20%, Nb: 3-5%, B: 0.001 to 0.02%, Ti: 0.3-1.3%, W: 5.5% or less, Mo: 2% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

11. An austenitic steel powder, comprising, in mass%, Ni: 30-40%, Cr: 15-20%, Nb: 3.5-4.5%, B: 0.001 to 0.02%, Ti: 0.5-1.1% or less, W: 5.5% or less, Zr: 0.3% or less, and the balance of Fe and inevitable impurities.

12. The austenitic steel powder according to any of claims 9 to 11, wherein the austenitic steel powder has an average grain size of 250 μm or less.