TIAL INTERMETALLIC CASTING ALLOY

Abstract

The present disclosure relates to the technical field of TiAl intermetallic foundry alloys. More particularly, it relates to such an alloy comprising: 44≤Al≤47 at. %; 0<Zr<2 at. %; 1≤W≤2 at. %; and only one of the following elements: 0.5≤V≤2 at. %; 0.5≤Cr≤2 at. %; 0<Nb<4 at. %.

Description

TECHNICAL FIELD OF THE INVENTION

The present invention relates to the technical field of TiAl intermetallic alloys and in particular foundry intermetallic alloys for manufacturing aeronautical turbines such as vanes or distributors. More particularly, the present invention relates to TiAl intermetallic alloys compatible with fast turbines making it possible to improve the thermodynamic efficiency of the engines and thus to limit polluting emissions.

PRIOR ART

Currently, the turbines most commonly used are slow turbines, such as for example CFM56, LEAP or GE90 turbines. The parts of these turbines, in particular the vanes and/or the distributors, are generally produced from first generation TiAl intermetallic alloy, such as for example Ti-48Al-2Cr-2Nb (in at. %—hereinafter TiAl 48-2-2). These alloys have mechanical properties which are sufficient for the temperatures reached by these turbines and the mechanical stresses to which they are subjected.

However, new turbine technologies are currently being studied, in particular for the purpose of improving the thermodynamic efficiency of the engines and thus limiting polluting emissions. One example of such turbines is the so-called “fast” turbines that are currently emerging, including GTF-1100G. The temperatures reached by these fast turbines are higher than by slow turbines (generally below 750° C.). The mechanical stresses that fast turbines must withstand are also greater.

TiAl 48-2-2 has its limits not because of the temperature of use, its mechanical properties being stable up to approximately 800° C., but because of the limits of acceptability of these. This is because they are judged to be too low for the requirements of fast turbines: with a conventional elastic limit Rp_0.2 (i.e. the stress leaving 0.2% residual plastic deformation after it is removed) of around 300 MPa over the temperature range of interest, in particular between 25° C. and 900° C. and an elongation at break close to 1% at ambient temperature.

Furthermore, this alloy undergoes peritectic solidification and usually has a duplex microstructure composed of γ grains (γ-TiAl) and lamellar grains (γ-TiAl and α₂-Ti₃Al lamellae).

Another alloy is then used: Ti-43.5Al-4Nb-1Mo-0.1B (at. %—hereinafter TNM-B1), which has mechanical properties superior to TiAl 48-2-2 over the temperature range of 25 to 750° C. It has a high Rp_0.2 greater than 500 MPa over the temperature range of 25° C. to 750° C. It does however have weaknesses. First of all, its elongation at break at ambient temperature is small (below 1%). Moreover, this alloy contains high quantities of β-genic alloy element (#phase stabilizer), such as niobium (Nb) and molybdenum (Mo). This leads to the formation of a triplex microstructure composed of γ grains (γ-TiAl), lamellar grains (γ-TiAl and α₂-Ti₃Al lamellae), and β grains (β-TiAl). Although β grains improve forging of this alloy at high temperature (above 1200° C.) and facilitate the production of parts to the required dimensions, these β grains, in the case of this alloy, remain in the microstructure at operating temperatures despite the applied heat treatments and which are aimed at eliminating them. Consequently, the mechanical properties of the alloy are greatly reduced: at temperatures above 750° C. the mechanical properties of TNM-B1 decrease very rapidly in traction and flow, and in certain cases can be inferior to those of TiAl 48-2-2. Moreover, this drawback goes also together with greater sensitivity to the phenomena of ageing in air, making the alloy even less ductile, a drawback that is not present with TiAl 48-2-2.

In order to improve the mechanical strength of these alloys when hot, some authors suggest combinations of addition elements for making the TiAl matrix stronger. For example, it has been suggested that tungsten (W) be added to strengthen the alloy.

Document U.S. Pat. No. 5,286,443 A describes a TiAl intermetallic alloy with the following formula:

Ti_xEI_γMe_zAl_1-(x+y+z),

EI representing B, Ge or Si; Me representing Co, Cr, Ge, Hf, Mn, Mo, Nb, Pd, Ta, V, W, Y, and/or Zr; 0.46≤x≤0.54; y being able to take several values depending on the selected EI element, but in general is between 0 and 0.02; 0.01≤z≤0.04 if a single Me element is selected, or 0.01≤z≤0.08 if a plurality of Me elements are selected; 0.46≤x+y+z≤0.54. No fewer than 77 examples of the composition are presented in this document. However, apart from hardness and breaking strength, which are discussed briefly, this document provides no information about the crystalline structure of the alloy. Now, as we have seen for TNM-B1, what is detrimental to the mechanical properties of the alloy is the presence of the β phase, and elements W, Nb and Ta are known for being β-genic and therefore favor the formation of β grains in the alloy. Moreover, all the alloys presented in this document comprise 46 and 54 at. % of aluminum whereas it is known that such a quantity of aluminum favors peripatetic solidification leading to severe segregation of the β-genic elements as well as a large grain size.

Presentation of the Invention

One of the objectives of the present invention is therefore to overcome at least one drawback of the prior art discussed above and more particularly to provide a TiAl intermetallic foundry alloy having both good mechanical properties at high temperature and sufficient ductility at ambient temperature.

Thus the present invention proposes a TiAl intermetallic foundry alloy comprising:

• • 44≤Al≤47 at. %; 0<Zr<2 at. %; 1≤W≤2 at. %; and only one of the following elements:
  • 0.5≤V≤2 at. %; 0.5≤Cr≤2 at. %; 0<Nb<4 at. %

Such an alloy has mechanical properties intermediate between TiAl 48-2-2 and TNM-B1, while having increased stability and resistance to oxidation at temperatures above 700° C.

Moreover, such a composition allows β solidification of the alloy. β solidification is particularly sought here since it makes it possible for the raw foundry grain size to become smaller and thus to reduce segregation risks and also improve ductility of the alloy. The β domains can then be eliminated by heat treatments.

Such an alloy can be used in manufacturing parts of aeronautical fast turbines, such as the blades and distributors.

Other optional and non-limiting features are as follows.

The alloy preferentially comprises less than 0.3 at. % of Si, C and B in total.

When the alloy comprises vanadium, it preferably comprises: 44≤Al≤46 at. %, preferably 44.5≤Al≤45.5 at. %; 0<Zr<2 at. %; 1≤W≤2 at. %; and 0.5≤V≤2 at. %. It is free from chromium and niobium.

When the alloy comprises chromium, it preferably comprises: 44≤Al≤46 at. %, preferably 44.5≤Al≤45.5 at. %; 0<Zr<2 at. %; 1≤W≤2 at. %; and 0.5≤Cr≤2 at. % It is free from vanadium and niobium.

When the alloy comprises niobium, it preferably comprises: 45≤Al≤47 at. %, preferably 45.5≤Al≤46.5 at. %; 0<Zr<2 at. %; 1≤W≤2 at. %; and 0<Nb<4 at. % It is free from vanadium and chromium.

The alloy may preferably comprise a β phase fraction of less than 1 vol. %.

The alloy may preferably have a single-phase domain at a temperature below 1350° C., preferably below 1300° C.

The alloy may have a conventional elastic limit Rp0.2 of between 400 and 700 MPa at a temperature between 25 and 900° C.

The alloy may comprise at least 20% α2 phase by volume at ambient temperature (in particular between 20° C. and 25° C.).

The present alloy is advantageously used for manufacturing aeronautical parts and in particular for fast turbines, for example blades or distributors.

BRIEF DESCRIPTION OF THE FIGURES

FIGS. 1 to 11 show photographs of the microstructure of the alloys of the examples and comparative examples. The contrast has not been homogenized over all the photographs.

FIG. 1 shows the microstructure straight from solidification of comparative example 1 (Ti-45Al-2W), β phase dendrites DB testifying to β solidification can be observed.

FIG. 2 shows the microstructure straight from solidification of comparative example 4 (Ti-45Al-2W-1Zr), β phase dendrites DB testifying to β solidification can be observed.

FIG. 3 shows the microstructure straight from solidification of comparative example 5 (Ti-45Al-2W-2Zr), β phase dendrites DB testifying to β solidification can be observed.

FIG. 4 shows the microstructure after TT1 of comparative example 1, a large β phase fraction PB can be observed.

FIG. 5 shows the microstructure after TT1 of comparative example 4, a large β phase fraction PB can be observed.

FIG. 6 shows the microstructure after TT1 of comparative example 5, a large β phase fraction PB can be observed, in places precipitated in lamellar form LB.

FIG. 7 shows the microstructure after TT1 of comparative example 6, (Ti-45Al-2W-1Zr-1V-3Nb), a large β phase fraction can be observed, precipitated in lamellar form LB.

FIG. 8 shows the microstructure after TT1 of comparative example 8, (Ti-45Al-2W-1Zr-1Cr-3Nb), a large β phase fraction can be observed, precipitated in lamellar form LB.

FIG. 9 shows the microstructure after TT1 of example 1 (Ti-45Al-2W-1Zr-1V) according to the invention, the almost total absence of β phase PB can be observed.

FIG. 10 shows the microstructure after TT1 of example 2 (Ti-45Al-2W-1Zr-1Cr) according to the invention, a small β phase fraction PB can be observed.

FIG. 11 shows the microstructure after TT1 of example 5 (Ti-46Al-2W-1Zr-3Nb) according to the invention, a small β phase fraction PB can be observed.

DETAILED DESCRIPTION OF THE INVENTION

The TiAl intermetallic foundry alloy according to the present invention is described further below.

This TiAl intermetallic foundry alloy comprises titanium, aluminum, zirconium, tungsten and a single element between vanadium, chromium and niobium. More particular, this foundry alloy comprises 44≤Al≤47 at. %; 0<Zr<2 at. %; 1≤W≤2 at. %; and 0.5≤Cr≤2 at. % or 0.5≤Cr≤2 at. % or 0<Nb<4 at. %.

The quantity of aluminum was selected subsequently to the following observations. The more the proportion of aluminum increases, the more the elastic limit Rp_0.2 at ambient temperature decreases, and the more the ductility increases. Thus, to be able to have mechanical performances lying between TiAl 48-2-2 and TNM-B1, it is necessary to maintain the aluminum between 44 and 48 at. %.

It is moreover known that tungsten improves mechanical strength and resistance to oxidation. Adding tungsten is however not obvious since this element is β-genic, i.e. the β phase fraction remaining despite heat treatments aimed at reducing it is great. However, in the context of the present invention, adding tungsten affords stabilization of the α single-phase domain at high temperature and enables it to be maintained over the mentioned range of aluminum contents. Below 1 at. % of tungsten, no beneficial effect of adding this element for the alloy is observed. Above 2 at. %, the advantage brought about by adding tungsten is counteracted by the appearance of the β phase in excessive quantities and by a degree of segregation too high to be eliminated through heat treatments or by adding alloy elements. The chosen range thus enables the alloy to preferably comprise less than 1 vol. % of β phase.

It has been observed by the authors that zirconium makes it possible to stabilize the lamellar microstructure of the alloy, with the presence of γ grains at the grain joints, but also to reduce segregation, including during casting, and to limit formation of β phase. This thus makes it possible to avoid deterioration of the alloy mechanical properties. However, it has been observed that, starting from 2 at. %, although foundry segregations are greatly reduced, on the other hand the β phase fraction increases beyond 1 vol. % and neither heat treatment nor adding addition elements makes it possible to go back below this limit.

Vanadium makes it possible to greatly limit segregations during casting and improve ductility of TiAl alloys at low temperatures. Vanadium can also increase the creep resistance of these alloys.

Chromium improves the resistance to oxidation and the ductility of TiAl alloys at low temperatures.

Niobium makes it possible to improve the mechanical strength and the resistance to oxidation of TiAl alloys. However, this element has a strong β-genic power that can cause precipitation of a large β phase fraction that does not disappear during heat treatment. Thus, it was not obvious for a person skilled in the art to add niobium in addition to W and Zr.

The combination of the proportions of aluminum, zirconium, tungsten and one of the elements mentioned among vanadium, chromium and niobium enables the alloy according to the invention to have solidification through the β phase which, as indicated above, makes it possible to refine the grain size straight from casting and thus reduce the risks of segregations as well as improve the ductility of the alloy while having the properties needed at high temperature, for example mechanical strength and resistance to oxidation. Although β solidification is sought since it leads to finer grains, it must be ensured that it is possible to control the residual β phase fraction over the temperature range between 25 and 900° C. This is because only part of the β phase is transformed into a phase during cooling. The combination mentioned above on the other hand makes it possible to eliminate a major part of the β phase by heat treatment until it drops below 1% by volume.

Preferably, the alloy comprises less than 0.3 at. % of Si, C and B in total, preferably less than 0.1 at. %. A larger quantity of Si, C and/or B favors the appearance of the β phase that it is not possible to remove by heat treatment.

The alloy preferably has a single-phase domain at a temperature below 1350° C., preferably below 1300° C. It preferably has a conventional elastic limit Rp_0.2 between 400 and 700 MPa at a temperature between 25 and 900° C. This means that, between 25 and 900° C., the conventional elastic limit remains between 400 and 700 MPa.

The alloy described above can be obtained by the following manufacturing method.

The alloy is cast from Ti-45Al-2W while adding the necessary addition elements, i.e.: zirconium and only a single one of the elements vanadium, chromium and niobium. The quantities to be added must make it possible to achieve the characteristics of the alloy described above.

The cast alloy is subjected to a homogenization treatment followed by controlled cooling.

The homogenization treatment comprises heating at a temperature of between 1200° C. and 1500° C., preferably between 1250° C. and 1350° C., for example 1300° C. The heating time at the selected temperature can be between 1 and 15 hr, preferably between 1 hr and 10 hr, for example 5 hr. The homogenization treatment is preferably implemented under neutral gas.

The homogenization treatment makes it possible to homogenize the solidification microstructure while maintaining the β phase fraction at a minimum value (in particular below 1 vol. %).

Cooling is performed in a controlled manner at a speed of between 25° C./min and 300° C./min, preferentially 50° C./min and 200° C./min, for example 100° C./min. Cooling is preferably applied until ambient temperature is reached. This cooling makes it possible to saturate the a phase at high temperatures so that, at the second ageing treatment described below, α₂+γ lamellar grains can be obtained with fine lamellae.

A second ageing treatment at a temperature below the homogenization temperature may be provided. This ageing comprises heating at a temperature of between 700° C. and 1200° C., preferably between 850° C. and 1100° C., for example 900° C. or 1000° C. Cooling is preferably performed under neutral gas. Ageing may last for 3 to 15 hr, preferably from 4 to 10 hr, for example 3 hr or 6 hr. Ageing is preferably performed under neutral gas. The temperature ranges mentioned above make it possible to minimize the discontinuous growth type transformation that is unfavorable to creep resistance. This second ageing treatment also makes it possible to relax internal stresses caused by cooling following the homogenization treatment.

EXAMPLES

In one example, the alloy is an alloy comprising vanadium but is free from chromium and niobium. This alloy comprises:

• • 44≤Al≤46 at. %, preferably 44.5≤Al≤45.5 at. %, for example 45 at %;
  • 0<Zr<2 at. %, preferably 0.5≤Zr≤1.5 at. %, for example 1 at %;
  • 1≤W≤2 at. %, for example 1 at. %, 1.5 at. % or 2 at. %; and
  • 0.5≤V≤2 at. %, for example 1 at. % or 1.5 at. %.

In one example, the alloy is an alloy comprising chromium but is free from vanadium and niobium. This alloy comprises:

• • 44≤Al≤46 at. %, preferably 44.5≤Al≤45.5 at. %, for example 45 at %;
  • 0<Zr<2 at. %, preferably 0.5≤Zr≤1.5 at. %, for example 1 at %;
  • 1≤W≤2 at. %, for example 1 at. %, 1.5 at. % or 2 at. %; and
  • 0.5≤Cr≤2 at. %, for example 1 at. % or 1.5 at. %.

In yet another example, the alloy is an alloy comprising niobium but is free from vanadium and chromium. This alloy comprises:

• • 45≤Al≤47 at. %, preferably 45.5≤Al≤46.5 at. %, for example 46 at %;
  • 0<Zr<2 at. %, preferably 0.5≤Zr≤1.5 at. %, for example 1 at %;
  • 1≤W≤2 at. %, for example 1 at. %, 1.5 at. % or 2 at. %;
  • 0≤Nb≤4 at. %, preferably 0.5≤Nb≤3.5 at. %, for example 1 at %, 2 at. % or 3 at. %.

The following table 1 presents a plurality of alloys according to the invention. The proportions are given as an atomic percent.

TABLE 1
Example	Ti	Al	W	Zr	V	Cr	Nb
1	51	45	2	1	1	—	—
2	51	45	2	1	—	1	—
3	51	46	2	1	—	—	1
4	50	46	2	1	—	—	2
5	49	46	2	1	—	—	3

These alloys were cast from Ti-45Al-2W while adding thereto the desired edition elements.

Firstly, quenching tests were performed for different homogenization treatment temperatures, sometimes supplemented by ageing at a lower temperature in order to optimize the homogenization temperature for the high-temperature treatment. These tests make it possible to know the stable thermodynamic state of the alloy at various temperatures (1400° C., 1300° C. and 1200° C.).

Following these tests, the microstructure, in terms of present phases, morphology of the grains and of the phases as well, as the fraction of the present phases, was determined for the alloys of the examples by image analysis. Part of these results is summarized in table 2.

The following conditions were used for the tests, the results of which are presented in table 2 (and table 5, “TH”): heating at 1400° C. for 17 hr under neutral gas and then cooling under neutral gas to ambient temperature then ageing at a temperature of 1300° C. for 5 hr under neutral gas and oil-bath quenching. These conditions correspond to an example of treatment that can be applied industrially to the alloy according to the invention.

	TABLE 2
			β phase
	Example	Microstructure	(vol. %)
	1	L + γ	0
	2	L + γ	0.1
	3	L + γ + β	0.7
	4	L + γ + β	0.7
	5	L	0

In table 2, “L” means lamellar grains, “γ” means γ grains, and “β” means β grains.

At the end of the quenching tests, optimized and industrially applicable heat treatments were determined. For example, the following conditions were determined (“TT” in table 5): heating at 1300° C. for 5 hr and then cooling under neutral gas to ambient temperature and heating at 900° C. for 6 hr and then furnace cooling (i.e. slow cooling due to the inertia of the furnace).

			β phase
	Example	Microstructure	(vol. %)
	1	L + γ + β	0.7
	2	L + γ + β	0.1
	3	L + γ + β	1.2
	4	L + γ + β	2.3
	5	L + β	0.2

In table 3, “L” means lamellar grains, “y” means γ grains, and “β” means β grains.

Comparative Examples

Comparative examples were performed in order to compare their properties with those of the above examples. Their compositions are provided in table 4. The proportions are given as atomic percentages.

TABLE 4
Comparative
example	Ti	Al	W	Zr	V	Cr	Nb
1	53	45	2	—	—	—	—
2	51	46	2	—	—	—	—
3	50	47	2	—	—	—	—
4	52	45	2	1	—	—	—
5	51	45	2	2	—	—	—
6	48	45	2	1	1	—	3
7	49.5	45	2	1	1	—	1.5
8	48	45	2	1	—	1	3
9	49.5	45	2	1	—	1	1.5
10	48	46	2	1	1	—	2
11	48	46	2	1	—	1	2

The same quenching tests TH and heat treatment tests TT described above were carried out. Some results are summarized in table 5.

TABLE 5
Comparative	Heat		B phase
example	treatment	Microstructure	(vol. %)
1	TH	L + β	0.1
2	TH	L + γ + β	0.3
3	TH	L	0
4	TT	L + β + γ	0.8
5	TT	L + β + γ	3.5
6	TT	L + β	19.3
7	TH	L + β + γ	6.4
7	TT	L + β + γ	4.5
8	TH	L + β	9
8	TT	L + β	8.6
9	TH	L + β + γ	5.1
9	TT	L + β + γ	5.5
10	TT	L + β + γ	10.5
11	TT	L + β + γ	15.4

In table 3, “L” means lamellar grains, “y” means γ grains, and “β” means β grains.

Discussion

The comparison between comparative examples 1, 4 and 5 (FIGS. 1 to 6) shows that adding at least 1 at. % zirconium makes it possible to stabilize the lamellar microstructure of the alloy, with the presence of γ grains at the grain joints, but also to reduce segregations. The reduction of segregations is obtained not only after the heat treatments (FIGS. 5 and 6) but also already at the time of casting (FIGS. 2 and 3). Moreover, the results show that this element limits the formation of β phase. However, beyond 2 at. %, although foundry segregations are greatly reduced, the β phase fraction increases substantially beyond the limit acceptable for the envisaged applications and precipitates in the form of lamellae (FIG. 6).

The results of comparative examples 6 to 11 (table 4) show that, instead of improving the mechanical properties and resistance to oxidation, combining at least two elements from vanadium, chromium and niobium leads to the formation of an excessively large β phase fraction, well beyond the limit acceptable for the envisaged applications (FIGS. 7 and 8). Thus, these elements have a disadvantageous synergistic effect.

On the other hand, the results of examples 1 to 5 show that the β phase fraction is limited to less than 1 vol. % after heat treatment TH (table 2, for all examples) and optimized heat treatments TT (table 3 for examples 1, 2 and 5; FIGS. 9 to 11).

Claims

1. A TiAl intermetallic foundry alloy comprising:

44≤Al≤47 at. %;

0<Zr<2 at. %;

1≤W≤2 at. %; and

only one of the following elements:

0.5≤V≤2 at. %;

0.5≤Cr≤2 at. %;

0<Nb<4 at. %.

2. The alloy according to claim 1, comprising less than 0.1 at. % of Si, C and B in total.

3. The alloy according to claim 1, comprising:

44≤Al≤46 at. %;

0<Zr<2 at. %;

1≤W≤2 at. %; and

0.5≤V≤2 at. %.

4. The alloy according to claim 1, comprising:

44≤Al≤46 at. %;

0<Zr<2 at. %;

1≤W≤2 at. %; and

0.5≤Cr≤2 at. %.

5. The alloy according to claim 1, comprising:

45≤Al≤47 at. %;

0<Zr<2 at. %;

1≤W≤2 at. %; and

0<Nb<4 at. %.

6. The alloy according to claim 1, comprising a β phase fraction of less than 1 vol. %.

7. The alloy according to claim 1, having a single-phase domain at a temperature below 1350° C.

8. The alloy according to claim 1, having a conventional elastic limit Rp0.2 of between 400 and 700 MPa at a temperature between 25 and 900° C.

9. The alloy according to claim 1, comprising at least 20% α2 phase by volume.

10. A part for a turbine produced in the alloy according to claim 1.

11. A method for manufacturing the alloy according to claim 1, comprising, in the following order:

casting a mixture of precursors of Ti, Al, Zr, W and only one of the elements V, Cr and Nb;

homogenizing at a temperature between 1200° C. and 1500° C., for a time between 1 hr and 15 hr;

optionally ageing at a temperature below the homogenization temperature at a temperature between 700° C. and 1200° C. for 3 hr to 15 hr; and

controlled cooling.

12. The method according to claim 11, wherein said homogenizing at a temperature of 1200° C. and 1500° C. includes homogenizing at a temperature of 1250° C. and 1350° C.

13. The method according to claim 11, wherein said homogenizing at a temperature of 1200° C. and 1500° C. for a time of between 1 hr and 15 hr includes homogenizing at a time between 1 hr and 10 hr.

14. The method according to claim 11, wherein said optionally ageing at a temperature below the homogenization temperature at a temperature between 700° C. and 1200° C. includes optionally ageing at a temperature between 850° C. and 1100° C.

15. The method according to claim 11, wherein said optionally ageing at a temperature below the homogenization temperature at a temperature between 700° C. and 1200° C. for 3 hr to 15 hr includes optionally ageing for a time between 4 hr to 10 hr.

16. The method according to claim 11, wherein said homogenizing and optionally ageing occurs under neutral gas pressure.

17. The alloy according to claim 3, wherein Al at. % is 44.5≤Al≤45.5 at. %.

18. The alloy according to claim 4, wherein Al at. % is 44.5≤Al≤45.5 at. %.

19. The alloy according to claim 5, wherein Al at. % is 44.5≤Al≤45.5 at. %.

20. The alloy according to claim 1, wherein the single-phase domain is at a temperature below 1300° C.