EP1420075B1 - Nickel-base superalloy - Google Patents

Description

Technical area

The invention relates to the field of materials technology. It relates to a nickel-base superalloy, in particular for the production of single-crystal components, such as blades for gas turbines.

State of the art

Such nickel-base superalloys are known. Single crystal components of these alloys have a very good material strength at high temperatures. As a result, z. B. the inlet temperature of gas turbines are increased, whereby the efficiency of the gas turbine increases.

Nickel-based superalloys for single-crystal components, as known from US Pat. Nos. 4,643,782, EP 0 208 645 and US Pat. No. 5,270,123, contain mixed-crystal-hardening alloying elements, for example Re, W, Mo, Co, Cr, and γ'-phase-forming elements, for example, Al, Ta, and Ti. The content of refractory alloying elements (W, Mo, Re) in the base matrix (austenitic γ-phase) continuously increases with the increase of the stress temperature of the Alloy. To contain z. B. common nickel-based superalloys for single crystals 6-8% W, up to 6% Re and up to 2% Mo (in% by weight). The alloys disclosed in the above references have high creep strength, good LCF (low duty cycle fatigue) and HCF (high cycle fatigue) properties, and high oxidation resistance.

These known alloys have been developed for aircraft turbines and therefore optimized for short and medium time use, i. the load duration is designed for up to 20,000 hours. In contrast, industrial gas turbine components have to be designed for a service life of up to 75,000 hours.

After a period of use of 300 hours z. For example, the alloy CMSX-4 of US 4,643,782 when experimentally used in a gas turbine at a temperature above 1000 ° C, a strong coarsening of the γ'-phase, which is associated with an increase in the creeping speed of the alloy adversely.

Also, for example, the alloys known from US 5,270,123 have comparable disadvantages. By the alloying elements selected there, a positive or a negative lattice offset between the matrix forming γ phase and the γ 'phase, ie the secondary intermetallic phase Ni ₃ Al, in the case of the Ta, Ti, Hf partial Al and Co, and Cr can partly replace Ni. This lattice distortion hinders dislocations when sliding or cutting the γ'-grains. Although the lattice distortion causes an increase in the Kurrzeitfestigkeit, but with prolonged stress, a coarsening of the microstructure and then a degradation of the γ 'structure and thus causes a long-term mechanical weakening of the alloy.

This disadvantage is eliminated with the alloy known from EP 0 914 483 B1. This nickel base superalloy consists essentially of (measured in% by weight) 6.0-6.8% Cr, 8.0-10.0% Co, 0.5-0.7% Mo, 6.2-6.6% W, 2.7-3.2% Re, 5.4- 5.8% Al, 0.5-0.9% Ti, 7.2-7.8% Ta, 0.15-0.3% Hf, 0.02-0.04% C, 40-100 ppm B, 0-400 ppm Y, balance Ni with impurities, the ratio of ( Ta + 1.5 Hf + 0.5 Mo - 0.5 Ti) / (W + 1.2 Re) ≥ 0.7. These alloys have no lattice offset between the γ-phase and the γ'-phase due to the said ratio of the alloying elements at the operating temperature, whereby a high long-term stability under moderate load is achieved. In addition, this nickel-base superalloy alloyed with rhenium has excellent castability and high phase stability combined with the best mechanical properties. It is also characterized by high fatigue strength and creep stability even with long-term exposure.

It has also been found that in the presence of mechanical stress and a long-term high temperature stress to a directional coarsening of the y'-particles, the so-called rafting (rafting) occurs, and, at high γ'-levels (ie at a γ'-volume fraction of at least 50%) for inversion of the microstructure, ie γ 'becomes the continuous phase in which the former γ-matrix is embedded. Since the intermetallic γ'-phase tends to environmental embrittlement, under certain loading conditions this leads to a massive decrease of the mechanical properties, especially the yield strength, at room temperature (degradation of the properties). The environmental embrittlement occurs especially when moisture and long hold times are under tensile load.

Presentation of the invention

The aim of the invention is to avoid the disadvantages mentioned. The invention is based on the object to develop a nickel-based superalloy, which on the one hand has a strong and strong γ-phase as a matrix and which on the other hand only a small proportion, i. less than 50%, at γ'-phase, and thus is very resistant to oxidation and has a good creep behavior.

• 7-13 Cr
• 4-10 Co
• 0.5-2 Mo
• 2-8 W
• 4-6 Ta
• 3-6 Al
• 1-4 Ti
• 0.1-6 Ru
• 0.01-0.5 Hf
• 0.001-0.15 Si
• 0-700 ppm C
• 0-300 ppm B

According to the invention, this object is achieved in that the nickel-based superalloy according to the invention is characterized by the following chemical composition (data in% by weight):

• 7-13 Cr
• 4-10 co
• 0.5-2 mo
• 2-8 W
• 4-6 d
• 3-6 Al
• 1-4 Ti
• 0.1-6 Ru
• 0.01-0.5 Hf
• 0.001-0.15 Si
• 0-700 ppm C
• 0-300 ppm B

Remaining nickel and manufacturing-related impurities.

The advantages of the invention are that the alloy has a good degradation behavior. The γ phase (matrix) is solidified by the addition of ruthenium, despite the absence of rhenium, which according to the known state of the art is considered to be a particularly good solid solution promoter and therefore greatly improves the properties of the γ matrix. The alloy according to the invention is distinguished by good creep rupture strength, stable microstructures and good castability.

In addition, the oxidation resistance of the alloy is very good. It is outstandingly suitable for the production of single-crystal components, for example blades for gas turbines.

Due to the small proportion of secondary precipitation-hardening γ'-phase, which is incorporated in the strongly solidified γ-phase, the degradation behavior of the alloy according to the invention is good. There is no single crystal crack growth and no large decrease in yield strength at room temperature in the degraded state compared to the non-degraded state.

• 10-13 Cr
• 8-9 Co
• 1.5-2 Mo
• 3-5 W
• 4-5 Ta
• 3-5 Al
• 2-4 Ti
• 0.3-4 Ru
• 0.01-0.5 Hf
• 0.001-0.15 Si
• 0-700 ppm C
• 0-300 ppm B

Preferred ranges of the nickel-base superalloy according to the invention are (in% by weight):

• 10-13 Cr
• 8-9 Co
• 1.5-2 Mo
• 3-5 W
• 4-5 d
• 3-5 al
• 2-4 ti
• 0.3-4 Ru
• 0.01-0.5 Hf
• 0.001-0.15 Si
• 0-700 ppm C
• 0-300 ppm B

Remaining nickel and manufacturing-related impurities.

• 10-13 Cr
• 8-9 Co
• 1.5-2 Mo
• 3.5-4 W
• 4-5 Ta
• 3.5-5 Al
• 3-4 Ti
• 0.3-1.5 Ru
• 0.5 Hf
• 10-500 ppm Si
• 250-350 ppm C
• 80-100 ppm B

A particularly preferred range of the nickel-based superalloy according to the invention is the following:

• 10-13 Cr
• 8-9 Co
• 1.5-2 Mo
• 3.5-4 W
• 4-5 d
• 3.5-5 Al
• 3-4 Ti
• 0.3-1.5 Ru
• 0.5 Hf
• 10-500 ppm Si
• 250-350 ppm C
• 80-100 ppm B

Remaining nickel and manufacturing-related impurities.

• 7-9 Cr
• 8-9 Co
• 1.5-2 Mo
• 3-5 W
• 5-6 Ta
• 3-5 Al
• 1-2 Ti
• 0.5-1.5 Ru
• 0.5 Hf
• 700 ppm C
• 100 ppm B
• 500 ppm Si

A further nickel-based superalloy according to the invention has the following chemical composition (data in% by weight):

• 7-9 cr
• 8-9 Co
• 1.5-2 Mo
• 3-5 W
• 5-6 d
• 3-5 al
• 1-2 Ti
• 0.5-1.5 Ru
• 0.5 Hf
• 700 ppm C
• 100 ppm B
• 500 ppm Si

Rest Nicckel and manufacturing impurities.

Brief description of the drawings

Fig. 1

ein Gefügebild der Vergleichslegierung VL;

Fig. 2

ein Gefügebild der erfindungsgemässen Legierung L1;

Fig. 3

ein Gefügebild der erfindungsgemässen Legierung L1 nach Degradierung;

Fig.4

ein Gefügebild der erfindungsgemässen Legierung L2 nach Degradierung;

Fig. 5

ein Diagramm, welches die Gewichtsänderung der Legierungen VL, L1 und L2 in Abhängigkeit von der Zeit angibt;

Fig. 6

ein Diagramm, welches die 0,2%-Streckgrenze der Legierungen VL, L1 und L2 in Abhängigkeit vom Degradations-Parameter angibt und

Fig. 7

ein Diagramm, welches die Spannung (1%-Dehngrenze) der Legierungen VL, L1 und L2 in Abhängigkeit vom Larson Miller-Parameter angibt.

In the drawings, two embodiments of the invention are shown. Show it:

Fig. 1

a micrograph of the comparative alloy VL;

Fig. 2

a micrograph of the inventive alloy L1;

Fig. 3

a micrograph of the inventive alloy L1 after degradation;

Figure 4

a micrograph of the inventive alloy L2 after degradation;

Fig. 5

a diagram indicating the weight change of the alloys VL, L1 and L2 as a function of time;

Fig. 6

a diagram indicating the 0.2% yield strength of the alloys VL, L1 and L2 depending on the degradation parameter, and

Fig. 7

a diagram which indicates the stress (1% proof stress) of the alloys VL, L1 and L2 in dependence on the Larson Miller parameter.

Ways to carry out the invention

The invention will be explained in more detail with reference to embodiments and FIGS. 1 to 7.

Nickel-based superalloys having the chemical composition given in Table 1 were investigated (in% by weight): Table 1: Chemical composition of the investigated alloys L1 (AMN1) L2 (AMN3) VL (PW 1483) Ni rest rest rest Cr 9.96 12:34 12.8 Co 8.86 8.84 9 Not a word 1:47 1.85 1.9 W 3:45 3.76 3.8 Ta 4 4.96 4 al 3:57 3:45 3.8 Ti 3.83 3.96 4 Hf 0.5 00:48 - C 0025 0033 - B 86 ppm 79 ppm - Si 10 ppm 10 ppm - Ru 1:07 00:28 -

Alloys L1 and L2 are alloys whose composition falls within the claims of the present invention. In contrast, alloy VL is a comparative alloy known in the art as PW 1483. It differs from the alloys according to the invention primarily in that it is not alloyed with ruthenium and no appreciable Si content is present. In the composition of the elements Cr, Co, Mo, Ta, Al, Ti and Ni, the alloys L2 and VL are almost identical. This is true up to the Cr content on the alloy L1. In the case of L1, the Cr content is about 3% by weight lower than in the case of the comparative alloy VL.

All three alloys were subjected to the following heat treatment: 1 h / 1204 ° C + 1 h / 1265 ° C + 4 h 1080 ° C.

The Vickers hardness HV2 was measured. The results listed in Table 2 were achieved. Table 2: Vickers hardness for the alloys tested L1 VL HV2 447 403

The alloy L1 thus has over 10% higher hardness than the comparative alloy VL. The γ-phase (matrix) of the alloys according to the invention is solidified mainly by the alloyed ruthenium.

FIG. 1 shows the microstructure of the comparative alloy VL1, while FIG. 2 shows the microstructure of the alloy L1 according to the invention.

In comparison with the alloy VL, the smaller proportion of γ'-phase (dark particles) is clearly recognizable in the alloy L1. The γ'-phase (secondary precipitation-hardening intermetallic phase) has an approximately quadrangular shape in the alloy VL and is arranged in a striped manner in the matrix. In contrast, the γ'-phase in L1 has a spherical shape, which is an indication of a very small lattice offset between the γ and the γ 'phase. This small lattice offset and, above all, the low volume fraction of γ'-phase (less than 50%) have a positive effect in that there is no γ / γ'-inversion of the microstructure, i. the γ'-phase is embedded in the γ-phase and does not form a continuous network. Thus, a good degradation behavior of the inventive alloys is achieved.

FIGS. 3 and 4 show micrographs of the novel alloys L1AD (FIG. 3) and L2AD (FIG. 4) in the degraded state (T = 1000 ° C., σ = 80 MPa, t = 747 h). The γ'-phase is embedded in the γ-phase and does not form a continuous network. The alloy L1AD shows predominantly round to oval forms of the γ'-phase, while in the alloy L2AD the γ'-phase is very elongated.

This has effects on the properties. In Fig. 5 the weight change as a function of time for the three alloys is shown. After being degraded, the alloys according to the invention have a significantly lower weight change than the comparative alloy known from the prior art, ie they have a significantly better oxidation resistance.

FIG. 6 shows the dependency of the 0.2% yield strength at room temperature on the degradation parameter P with $$\mathrm{P}=(\mathrm{T}-800){\mathrm{t}}^{1/2}{\mathrm{\sigma }}^{1/5}\mathrm{,}$$

While the comparative alloy VL and the alloy L2AD behave almost the same, for L1AD the voltage is about 200 MPa below the values for VL and L2AD.

auf, so ergeben sich die in Fig. 7 dargestellten Abhängigkeiten. Die Legierung L2AD weist über den gesamten Bereich höhere Dehngrenzen auf als die Vergleichslegierung (bei besserem Oxidationsverhalten). Zwar weist die Legierung L1AD nur geringere Dehngrenzen als die Vergleichslegierung VL auf, hat aber dafür ebenfalls eine wesentlich bessere Oxidationsbeständigkeit.If one carries the 0.1 proof stress over the Larson Miller parameter LM, with $$\mathrm{LM}=\mathrm{T}(\mathrm{log\; t}+20)$$

on, the dependencies shown in Fig. 7 result. The alloy L2AD has higher yield strengths over the entire range than the comparison alloy (with better oxidation behavior). Although the alloy L1AD has only lower yield strengths than the comparative alloy VL, but also has a much better oxidation resistance.

Of course, the invention is not limited to the described embodiments.

Claims (4)

1. Nickel-base superalloy for producing single-crystal components, characterized by the following chemical composition (details in % by weight):

   7-13 Cr

   4-10 Co

   0.5-2 Mo

   2-8 W

   4-6 Ta

   3-6 Al

   1-4 Ti

   0.1-6 Ru

   0.01-0.5 Hf

   0.001-0.15 Si

   0-700 ppm C

   0-300 ppm B

   
   remainder nickel and production-related impurities.
2. Nickel-base superalloy according to Claim 1, characterized by the following chemical composition (details in % by weight):

   10-13 Cr

   8-9 Co

   1.5-2 Mo

   3-5 W

   4-5 Ta

   3-5 Al

   2-4 Ti

   0.3-4 Ru

   0.01-0.5 Hf

   0.001-0.15 Si

   0.700 ppm C

   0-300 ppm B

   
   remainder nickel and production-related impurities.
3. Nickel-base superalloy according to Claim 2, characterized by the following chemical composition (details in % by weight):

   10-13 Cr

   8-9 Co

   1.5-2 Mo

   3.5-4 W

   4-5 Ta

   3.5-5 Al

   3-4 Ti

   0.3-1.5 Ru

   0.5 Hf

   10-500 ppm Si

   250-350 ppm C

   80-100 ppm B

   
   remainder nickel and production-related impurities.
4. Nickel-base superalloy according to Claim 1, characterized by the following chemical composition (details in % by weight):

   7-9 Cr

   8-9 Co

   1.5-2 Mo

   3-5 W

   5-6 Ta

   3-5 Al

   1-2 Ti

   0.5-1.5 Ru

   0.5 Hf

   500 ppm Si

   700 ppm C

   100 ppm B

   
   remainder nickel and production-related impurities.

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