KR100264089B1 - Oxidation resistant low expansion superalloy - Google Patents

Abstract

In the present invention, by weight ratio 26-50% cobalt, 20-40% nickel, 20-35% iron, 4-10% aluminum, 0.5-5% total niobium and 1/2 tantalum, 1.5-10% chromium, 0- 1% Titanium, 0-0.2% Carbon, 0-1% Copper, 0-2% Manganese, 0-2% Silicon, 0-8% Molybdenum, 0-8% Tungsten, 0-0.3% Boron, 0-2% Hafnium, 0-2% rhenium, 0-0.3% zirconium, 0-0.5% nitrogen, 0-1% yttrium, 0-1% lanthanum, 0-1% total rare earth metal except lanthanum, 0-1 It also contains% cerium, 0-1% magnesium, 0-1% calcium, 0-4% oxidative dispersion and impurities. The alloy is annealed at a temperature below 1010 ° C. or at a temperature between 1066 ° C. or 1110 ° C. and the melting point and also aged at beta penetration temperatures of 788 ° C. or higher to optimize crack resistance.

Alloy with controlled thermal expansion coefficient.

Description

Controlled thermal expansion superalloy

FIG. 1 is a graph showing static crack growth at 538 ° C. measured transversely- longitudinally for each composition.

FIG. 2 is a graph showing static crack growth at 538 ° C. temperature measured transversely to longitudinally to show the effect of Ni on crack growth rate. Pieces 6, 12 and 16 were annealed at 1010 ° C. for 1 hour, then air cooled and second heat treated at 788 ° C. for 16 hours, the furnace was cooled to 621 ° C., heat treated at this temperature for 8 hours and air cooled. do.

3 is the stress intensity 33 This is a static crack growth graph of 538 ° C. for an alloy annealed at 982 ° C. with different chromium contents in the transverse-longitudinal direction. The alloy was annealed at 982 ° C. for 1 hour, cooled to 621 ° C. for 8 hours, and then air cooled again.

4 is the stress intensity 33 Static crack growth graph of 538 ° C. of annealing and second heat treated alloys at different temperatures in the transverse-longitudinal direction at. The alloy was air cooled after annealing for 1 hour. Heat-treated for 16 hours at the temperature indicated in the figure, the furnace is cooled and left for 8 hours at 621 ℃ and then air cooled.

5 is the stress intensity 33 Is a graph showing the effect of chromium and cobalt content on the static crack growth rate at 538 ° C of samples tested transversely to longitudinally at.

FIG. 6 is a graph showing the effect of annealing temperature on da / dt as a function of Ni content for a second heat treated material below 1450 ° F. (788 ° C.).

7 is a graph showing the relationship between da / dt, crack plane direction, second creep velocity, annealing temperature and morphological aspects.

8 and 8 (a) show an annealing for one hour at the temperature shown and a second heat treatment for 16 hours at the temperature shown and k = 30 for an air cooled sample held at 1150 ° F. (621 ° C.) for 8 hours. (33 ) Three-dimensional graph showing the effect of annealing and second heat treatment on the 1000 ° F. (538 ° C.) crack growth rate for alloys with 27 to 32% nickel content at stress intensity.

FIG. 9 is a time-temperature-conversion graph for the hot-rolled piece 30 (Table 3) treated with 2100 ° C. (1149 ° C.) solution for 1 hour.

FIG. 10 is the complete da / dt crack growth curve at 538 ° C. for hot strips 30 (Table 3) tested in long and short longitudinal directions compared to alloy 718,909 and similar chromium free alloys.

The present invention relates to the field of thermal expansion controlled alloys. In particular, the present invention relates to the three phase, gamma, gamma prime, beta superalloy fields having relatively low coefficients of thermal expansion.

Three-phase low coefficient of thermal expansion alloys are known (EPO Patent Publication No. 433,072 ('072) published June 19, 1991). '072 offers improved resistance to grain boundary oxygen embrittlement (SAGBO) accelerated by low coefficient of thermal expansion and stress. The alloy of '072 has excellent notch burst strength, low density, and acceptable impact strength. Specific applications of Alloy 072 include turbine engine main components such as seals, rings, discs, compressor blades and casings. Low coefficient of thermal expansion alloys should not suddenly fail and are used in the field of structural components with close tolerances.

In the past, turbine engine manufacturers required alloys with only notch ductility for use in critical structural applications. Currently, turbine engine manufacturers need alloys that are resistant to crack growth. INCONEL Alloy 718 (registered trademark of Inco Alloys) is an example of a turbine alloy with excellent resistance to crack growth. Resistance to crack growth provides an alloy free of defects, voids and cracks. Moreover, resistance to crack growth makes it possible to predict crack location and component life by pre-failure inspection. Unfortunately, low thermal expansion coefficient superalloys to be used in combination with alloy 718 have crack growth problems at temperatures of 538 ° C. (1000 ° F.). Even though the '072 alloy provides excellent notch ductility and excellent resistance to crack initiation, more alloys with improved resistance to crack growth are needed.

INCOLOY alloy 909 (Incoalloy®) is used in applications where a low coefficient of thermal expansion is required. Low coefficient of thermal expansion (CTE) is defined as an alloy that provides at least 10% lower CTE than alloy 718. However, alloy 909 has a low coefficient of thermal expansion but does not provide resistance to crack growth of alloy 718. Moreover, alloy 909 causes extensive oxidation problems at elevated temperatures. Turbine engine components made from 909 and other 900 series alloys should be replaced periodically during engine maintenance. The replacement of components made from the 909 alloy accounts for a large proportion of the turbine engine's overall maintenance costs. Alloys with low thermal expansion along with oxidation resistance can reduce engine repair costs.

It is an object of the present invention to provide an alloy with improved resistance to crack growth in combination with SAGBO resistance, controlled low thermal expansion coefficient, notch rupture strength, impact strength and reduced density.

Another object of the present invention is to provide an alloy with low thermal expansion in combination with improved oxidation resistance and stability.

In the present invention, by weight ratio 26-50% cobalt, 20-40% nickel, 20-35% iron, 4-10% aluminum, 0.5-5% niobium and 1/2 weight ratio tantalum, also 1.5-10% chromium It provides an alloy of controlled low thermal expansion coefficient containing. In addition, this alloy contains 0-1% titanium, 0-0.2% carbon, 0-1% copper, 0-2% manganese, 0-2% silicon, 0-8% molybdenum, 0-8% tungsten, 0-0.3 0-1% total rare earth metals other than boron, 0-2% hafnium, 0-2% rhenium, 0-0.3% zirconium, 0-0.5% nitrogen, 0-1% yttrium, 0-1% lanthanum, lanthanum, Provided is a controlled coefficient of thermal expansion alloy containing 0-4% cerium, 0-1% magnesium, 0-1% calcium, 0-4% oxide dispersion and incidental impurities. The alloy may be annealed at a temperature below 1010 ° C. or at 1066 ° C. or 1110 ° C. to a melt temperature and further subjected to a second heat treatment at a beta settling temperature of 788 ° C. or higher to optimize resistance to crack growth.

Small amounts of chromium in combination with increased cobalt concentrations have been found to reduce the crack propagation rate. Furthermore, a four step heat treatment consisting of annealing, beta aging and two gamma frame aging steps was used in the presence of chromium to optimize crack growth and yield strength. In addition, this alloy produces a CTE reduction of more than 10% over the operating temperature range compared to alloy 718.

26% -50% cobalt increases resistance to crack growth at temperatures of 538 ° C. (unless otherwise noted, all compositions are expressed as weight ratios). Cobalt more than 50% lowers the burst strength. 20-40% nickel stabilizes the austenite phase. Moreover, nickel increases the room temperature ductility of the alloy. 20-35% iron provides a low coefficient of thermal expansion and lowers the bending temperature when replacing cobalt or nickel. Excess iron results in instability of the alloy.

Aluminum promotes beta phase formation. For the purposes of this application, the beta phase includes Al-rich phases that can be aligned and converted to intermetallic structures based on Al-lean FeAl, CoAl and NiAl. The beta phase is disordered at room temperature or higher. Beta phase alignment cooled to room temperature differs from beta alignment that occurs with high temperature applications. The beta phase contributes to providing resistance to stress accelerated grain boundary oxidation (SAGBO). The beta phase also contributes to the high temperature workability of the alloy. In addition, aluminum also promotes the formation of gammaprime phases that increase strength. It is believed that the beta phase and the gamma prime phase partially control the crack growth rate at 538 ° C. Finally, aluminum reduces the density of the alloy and also greatly improves surface oxidation resistance.

Containing a relatively small amount of 1.5-10% chromium in combination with a high content of cobalt increases the resistance to crack growth at high temperatures. Chromium also improves the response to heat treatment and increases the stress rupture strength. Preferably, 1.5-5% chromium is used resulting in a slight increase in CTE and a slight decrease in flexion temperature above the flexion temperature. Chromium also improves the creep resistance of the alloy.

0.5-5% niobium increases the tensile strength at high temperatures and also the high temperature stress rupture strength. Niobium also stabilizes the shape of the alloy and strengthens the beta phase.

Titanium up to 1% increases alloy strength. Excess titanium, however, promotes phase instability. The amount of carbon added is at most 0.2%. The increased carbon slightly reduces the stress rupture strength.

Copper is up to 1% and manganese is up to 2%. It is desirable to keep the silicon at 2% or less. Silicon has been shown to reduce stress rupture strength when less than 0.25%. Molybdenum up to 8% favors strength and increases corrosion resistance. However, molybdenum increases the density and the coefficient of thermal expansion. Tungsten up to 8% has been found to increase stress burst strength at the expense of density and coefficient of thermal expansion.

Boron may be present up to 0.3%. Excess boron causes high temperature malleability and weldability problems. Both hafnium and rhenium can be present in amounts up to 2%. Zirconium is present up to 0.3%. Zirconium can adversely affect high temperature malleability. Yttrium, tantalum and cerium may be present in amounts up to 1%. Other rare earth metals may be present up to 1%. Yttrium, lanthanum, cerium and rare earth metals are expected to increase oxidation resistance. It may contain up to 1% magnesium, calcium and other deoxidants and emollients. Alternatively, oxide dispersions such as yttria, alumina or zirconia can be used in amounts up to 4%. Preferably, the oxide dispersion is added by a mechanical alloy forming method.

Table 1 shows the composition of the present invention. Table 1 shows all ranges between the two specified values. For example, the alloy may comprise 28-40% Co, 25-30% Ni, 4.5-6% Al, 0.75-3.5% Nb and 1.5-5% Cr.

TABLE 1

Table 2 below shows the advantageous range of the invention believed to provide excellent resistance to crack growth at 538 ° C.

TABLE 2

aCu + Mn + Si ≤ 1.5

bMo + W ≤ 5

cCu + Mn + Si ≤ 1

Table 3 contains a list of compositions tested in the alloy of the present invention.

Table 4 shows the numbers of the segments which are listed sequentially in the composition of Table 3. All compositions in this specification are expressed in weight percent unless otherwise noted. Table 4 shows the lobes containing various amounts of nickel, cobalt, chromium and niobium, with 27.5% iron and 5.4% aluminum.

TABLE 3

TABLE 3a

TABLE 4

AC = air cooling

FC = furnace cooling

Table 5 provides the room temperature mechanical properties of the various alloys shown in Table 4.

TABLE 5

Table 5 shows that the strength and ductility of all materials containing 3% niobium satisfy gas turbine engine applications. The minimum at room temperature is 690 MPa (100 Ksi) 0.2% yield strength and the minimum at room temperature ductility is 10% elongation. More preferred is 0.2% yield strength at room temperature of about 825 MPa (120 Ksi). At 4% niobium, the strength of the alloy increases, but the ductility decreases. Chromium does not have a significant effect on strength, and ductility decreases if it contains more than 3,5% chromium.

Table 6 shows the mechanical properties of the alloy of Table 4 provided at 704 ° C.

TABLE 6

At elevated temperatures, the strength and ductility of all alloys are acceptable. Normally, the minimum elevated temperature is 590 MPa (85 Ksi) 0.2% yield strength (704 ° C) and the elevated temperature ductility minimum is 15% elongation (704 ° C). Increasing the nickel content has a great effect on the tensile strength at elevated temperatures. In general, chromium and niobium are somewhat advantageous in temperature rising properties.

Table 7 provides the effect of Cr-Nb-Ni on creep at elevated temperatures (ASTM E-139).

TABLE 7

Adding 2% chromium may increase the time to reach 0.2% strain by 100-400% compared to chromium-free alloys. In addition, the second creep rate is reduced several times in materials containing more than 2% chromium. Increasing nickel and niobium have a synergistic effect on creep properties. Adding 4% niobium to the 33% nickel-containing material lengthens the 0.2% strain time and reduces the second creep rate. Preferred creep parameters are at least 15 hours in 0.2% strain and the second creep speed is below 5 × 10 ⁻⁵ m / hr. Table 8 shows the effect of chromium-niobium and nickel on Charpy V-notch impact energy.

TABLE 8

The room temperature impact energy is low but suitable for turbine applications. The impact energy is equivalent to INCOLOY alloy 909. INCOLOY Alloy 909 has been successfully used in turbine manufacturing. Increasing nickel has been found to increase impact energy. The chromium effect is not critical and 4% niobium significantly reduces the impact energy. It is preferred to have a room temperature CVN impact energy of at least 5 Nm, in particular at least 10 Mm.

Table 9 shows the effect of chromium, nickel and niobium on the coefficient of thermal expansion at various temperatures.

TABLE 9

Below the bending temperature, the CTE decreases by 0.9 μm / m / ° C. when 0-2% chromium is added. At temperatures above the flexural temperature, the alloy shows an increase in CTE with ferromagnetic action. 2-4% chromium has little effect on the coefficient of thermal expansion in the ferromagnetic range below the bending temperature. However, chromium has a significant effect on CTE above flexural temperatures. Cobalt, however, tends to increase the bending temperature.

The CTE of the alloy is at least 10% lower than alloy 718 and less than 13.6 μm / m / ° C. at 649 ° C. It is desirable in many gas turbine designs to match the bending temperature and slope of INCONEL alloy 718 with a 10% reduction in CTE in the alloy of the present invention. For 4% chromium-containing alloys, CTE is 25% lower at 316 ° C and 21% lower at 427 ° C and 13% lower at 649 ° C. For 3% chromium-containing alloys, the CTE is 26% lower at 316 ° C, 23% lower at 427 ° C, and 16% lower at 649 ° C. Although the slope does not exactly match the slope of INCONEL alloy 718, the slope is sufficient to provide process advantages when using the alloy of the present invention in combination with alloy 718. Due to the addition of 4% chromium, even alloys with low bending temperatures have a suitable bending temperature for gas turbine engines. At temperatures above the flexural temperature, the coefficient of thermal expansion increases significantly.

In order to predict CTE for various Ni, Co and Cr weight ratio combinations, a linear regression model is formulated to correlate CTE at 316 ° C for alloys containing 27 Fe, 5.5Al and 3 Nb. The unit of this model is μm / m / ° C. and is as follows:

CTE _{316 ° C} = 3.64 + 0.007 (Co) (Ni)-0.281 (Cr) + 0.045 (Cr)

CTE _{649 ° C} = 12.58 + 0.099 (Cr) + 0.047 (Cr)-0.022 (Co)

It is experimentally found that the predictability of the above formula is excellent for 24-28% iron. Depending on the nickel content, the alloy contains up to 37% cobalt and up to 10% chromium and the CTE is kept 10% lower than alloy 718.

The 649 ° C model limits the maximum chromium content for best operation at elevated temperatures of up to 5, 5.5 and 6% chromium, depending on the cobalt concentration. In cases where the bending temperature is not exceeded, increasing the amount of chromium gives the desired CTE rate. Table 10 shows the effect of small amounts of chromium on corrosion resistance.

TABLE 10

Note: 1. Overall composition, see Table 3.

2. Conduct salt spray test by exposing 720 hours at 35 ℃ according to ASTM B117-85

3% chromium-containing materials have been found to eliminate corrosion that occurs in salt spray experiments according to ASTM B117-85. However, addition of only 1% chromium accelerates the point corrosion. The corrosion rate of the 3% chromium-containing material is superior to that of the INCOLOY alloy 909 compared to the 1% chromium-containing alloy. Molybdenum can replace some or all of chromium for salt spray resistance.

Table 11 contains the effects of chromium, niobium and nickel on static crack life (at 538 ° C).

TABLE 11

PCF = pre-crack

At 538 ° C, alloys such as INCOLOY alloys 907 and 909 are highly susceptible to cracking.

Break time or compressive tensile load crack life is improved by 1-2 times. The increase in crack life is particularly significant for alloys of low nickel concentrations and increased cobalt concentrations. Niobium appears to provide somewhat negative or no effect in high nickel alloys. Pre-cracks of 4% niobium and 27% nickel containing alloys show brittleness at room temperature. The alloy of the present invention had an initial stress strength of 27 And crack life for 10 hours at 538 ° C. More preferably, the crack life is 20 hours under the above conditions.

Table 12 shows the effect of chromium, niobium and nickel on static crack growth rate at 538 ° C.

TABLE 12

VT = Canceled Test

Table 12 shows that the static crack growth rate of alloys containing more than 2% chromium is reduced by 1-2 times. Alloys containing less than 30% nickel are particularly resistant to crack growth. The crack growth rate of the alloy containing 27% nickel is equal to the crack growth rate of conventional heat treated alloy 718. In FIG. 1, the resistance to crack growth of the alloy is improved 1 to 2 times by containing 2% or more of chromium. The alloy of ´072 has fewer defects or damage limits than are needed in certain structural applications. The alloy of the present invention containing at least 2% chromium is at the level of alloy 718. In fact, 50 At stress levels up to some alloys are more resistant to crack growth than alloy 718.

In particular, FIG. 2 shows the beneficial effect of reducing nickel and increasing cobalt on resistance to crack growth. Reducing nickel from 33% to 27% and cobalt from 28% to 34% improves resistance to crack growth. Particularly, cracks 16 containing 2.9% Cr, 27% Ni, 34% Co, and 28% Fe have good resistance to crack growth.

Table 13 shows the chromium-free alloy of '072 for comparison.

TABLE 13

PCF = pre-crack

The composition in Table 13 contains 33Ni-31Co-27Fe-5.3Al-3.0N with 0.02 chromium by weight percentage. The crack growth rate of the alloys in Table 11 is much higher than that of alloy 718. Heat treatment does not significantly affect the crack growth rate.

Table 14 shows the effects of various heat treatments on the static crack growth rate at 538 ° C.

TABLE 14

Stress intensity

Other heat treatment: 1010 ℃

788 ° C 33 4.2 × 10 ^-5

55 4.2 × 10 ^-4

1066 ℃

899 ℃ / 4 * 33 2.1 × 10 ^-5

55 2.5 × 10 ^-4

Alloy 718 33 1.3 × 10 ^-5

55 4.2 × 10 ^-5

* 899 = / 4h FC (38 ℃ / h) to 621 ℃ / 8h, AC

The composition in Table 14 contains 34 Ni-30 Co-24 Fe-5.4 Al-3.1 Cr-3.0 Nb in weight percent. In contrast to the alloys of Table 13, 3% chromium-containing alloys are positively affected by heat treatment. In FIG. 3, the crack growth rate of the alloy of the present invention subjected to annealing and aging is improved to match that of alloy 718. The alloy crack growth rate of ´072 is so large that it cannot be sufficiently improved by heat treatment.

The alloy of the present invention has a three phase structure. The main matrix is austenitic cubic or gamma phase. Beta phase provides SAGBO resistance. The gamma phase is enhanced by the gamma prime phase precipitation. In FIG. 4, the resistance to crack growth after high annealing temperature is improved by increasing the aging temperature and by β phase precipitation heat treatment. The beta phase is formed at annealing temperatures of less than 1090 ° C. (2000 ° F.). Beta phases are formed more at 750-1000 ° C (1382-1832 ° F). High temperature aging is particularly useful after high temperature soldering. Beta phase precipitation heat treatment is believed to contribute to reduced crack growth rate. Combining aging and cooling paths, eg cooling between furnace treatments at different temperatures, can primarily control the shape of the gamma prime strengthening phase.

Table 15 shows the effect of Cr, Ni annealing and aging on crack growth rate.

TABLE 15

caution

1) The da / dt rate in the 718 da / dt band is represented by a thick line.

2) Annealed: / 1h, AC

3) Aging: / 16h, furnace cooling 621 ℃ / 8h, AC

4) da / dt data from curved da / dt versus stress intensity curve

5) The test piece is a crimped tensile test piece, 7.62 mm thick × 24.4 mm width, precracked to 1.27 mm depth according to ASTM E647.

6) VT = Canceled Test

The data in Table 15 confirm the positive effects of chromium on crack growth rates. Furthermore, the reduction in nickel content leads to a decrease in crack growth rate. In addition to the composition change, the annealing temperature and aging temperature can also be adjusted to increase the resistance to crack growth. The crack growth aspect of the alloy of the present invention depends on the shape, volume percentage, position, etc. of the phase infiltrated into the alloy. A smaller proportion of the spherical beta phase is needed when the precipitate is at the grain boundaries. Beta alignment and conversion can play a role in resistance to crack growth.

In FIG. 5, cobalt concentration and chromium concentration greatly influence the crack growth rate. The data in FIG. 5 is based on alloys containing 24.5 to 27.5% Fe and 27 to 34% Ni. All alloys are annealed at 1010 ° C. for 1 hour, air cooled, aged at 778 ° C. for 16 hours and air cooled. 7 shows that the combination of trace amounts of chromium and high cobalt can improve resistance to crack growth. Preferably, the alloy of the present invention has a stress strength of 33 In addition, it is more preferable if it is 1x10 <^-4> mm / s or less at the temperature of 538 degreeC.

In FIG. 6, the nickel content reduction lowers the crack growth rate of the gamma-prime precipitation heat treated alloy. Maximum growth rates occur after 1900 ° F. (1038 ° C.) to 2000 ° F. (1093 ° C.) annealing temperatures. Minimum growth rates occur at annealing temperatures around 1800 ° F (982 ° C) or 2050 ° F (1121 ° C).

The Ni effect is very high, but especially when the material is annealed at 1900 ° F to 2000 ° F (1038 to 1093 ° C). Nickel content below 27% provides excellent da / dt resistance and resistance to crack initiation. The cracks containing 24% show an upper crack stop, which impairs the crack growth rate measurement force (the graph in FIG. 6 is the maximum crack growth rate that cannot explain why the crack growth was stopped during the test). However, 24% Ni-containing alloys have reduced stability, RTT strength and ductility, and at high ductility, stress rupture life is reduced. However, for 24% Ni-containing alloys this reduction in mechanical properties can be tolerated in some applications. Moreover, in some applications more than 24% nickel alloys are recommended for optimal combination of properties.

The annealing temperature and the Ni content and da / dt correlations are for aging heat treatments that do not contribute to da / dt resistance. Therefore, this graph shows that the optimum nickel content is between 26% and 29% at 1900 ° F annealing and the maximum Ni content when 1800 ° F (982 ° C) or 2050 ° F (1121 ° C) annealing and low temperature aging treatment are performed. Up to 34%.

Reducing the Co content and increasing Ni stabilizes the gamma prime by reducing the beta phase, or in some cases by changing the beta phase structure and composition to increase creep resistance, or aid in grain boundary oxygen diffusion, or both. Can be.

The flake 30 is obtained from about 4,000 kg vacuum induction application and vacuum arc remelting here. In FIG. 7, an engine ring of ten pieces 30 of 2 ″ (5.08 cm) thick × 4 ″ (10.16 cm) high × 28 ″ (71.12 cm) OD was tested, annealed, 1400 ° F. (760 ° C.) Aged for 12 hours at 1150 DEG F. (621 DEG C.), followed by air cooling for 8 hours.

The second creep rate decreases with increasing annealing temperature and is up to 1950 ° F (1066 ° C) for creep resistant superalloys. The creep rate decreases and the da / dt rate increases in the long transverse plane as expected. However, the da / dt in the short transverse plane did not change until the annealing temperature exceeded 1950 ° F. (1066 ° C.) and at this temperature the creep rate was increased to equal the da / dt in the long transverse plane.

After reaching the minimum value with an 1950 ° F (1066 ° C) annealing, the creep rate increases at 2000 ° F (1093 ° C) and 2050 ° F (1139 ° C) annealing. The long transverse plane da / dt decreases at the same annealing. Short transverse da / dt also decreases at 2050 ° F. (1139 ° C.) annealing.

This characteristic aspect is different from most superalloys treated with elevated temperature solutions. In general, creep speed decreases with increasing annealing temperature, resulting in coarse grain size. Superalloys in which cracks are grown by the environment have higher crack growth rates and rougher grain sizes.

The relationship between the annealing temperature of the da / dt and the creep rate is explained in part as a microstructural change. Four types of microstructures are distinguished when the annealing temperature increases.

After low temperature annealing below 1850 ° F. (1010 ° C.) (Type I) microstructures comprise fine grains and fine coarse beta-like particles in a “aggregate” structure with grain boundary precipitates. Most coarse beta particles are precipitated during pretreatment. Beta particles are softer than matrices at high processing temperatures, and beta particles before and during the treatment process are anisotropic. Fine grains and abundant beta degrade creep resistance and coarse creep rates. Gamma- during cold aging heat treatment (<1450 ° F or 788 ° C) by larger creep plasticity that blunts crack ends and grain boundary deposits that slow oxygen diffusion and longer crack paths (because of finer grains and coarse beta anisotropy). Frame settling occurs and da / dt is lowered.

As the annealing temperature increases in species II, the grain boundary beta precipitated during treatment begins to dissolve and the grains are granulated. Beta in granulated grains maintains anisotropy in species II. The small grain boundary beta increases the long axis da / dt, resulting in slower oxygen diffusion and better cracking paths due to the granulated grain. When grains are granulated and the beta content decreases, creep rate decreases. However, the uniaxial da / dt is unchanged or less because the crack plane must cross and pass through the longer beta particles. Beta particles blunt cracks or redistribute crack end stress and strain (because of local micro-creep plasticity).

Maximum long-term crack growth rate and minimum creep resistance occur at 1950 ° F (1060 ° C) annealing. At this annealing, the grain boundary precipitates are very small and the particle size is granulated to ASTM # 6 to # 4 (46 μm to 89 μm) but long beta is still present in the granulated grain (some of which are at the grain boundary triple point). .

Type III occurs at annealing temperatures of at least 1950 ° F. (1066 ° C.). The beta phase is reduced and the remaining beta particles are isotropic. There is little sediment between grains.

Grain size is slightly granulated and isotropic compared to 1950 ° F. annealed material. The uniaxial crack growth rate is greater than the long axial crack growth rate because there is no long beta to reduce crack growth along this direction.

However, the long axis da / dt is slightly lower and the creep speed is slightly higher. Therefore, submicron beta is precipitated or the gamma-prime structure is changed. Modifications of the beta phase atomic arrangement alter the creep mechanism.

In species IV, after annealing at 2050 ° F. (1121 ° C.), beta re-precipitation begins within the grain and at the boundaries. This precipitation may occur during the 1400 ° F. (760 ° C.) aging heat treatment cycle or during cooling from the 2050 ° F. (1121 ° C.) annealing temperature. Compared to the beta precipitated during thermomechanical treatment, this beta becomes very fine discrete particles at the grain boundaries and has a fine lath inside the grains. Beta particle reproducibility increases creep rate slightly and decreases long and short crack growth rates.

8 shows the effects of annealing and aging temperature on 538 ° C DA / DT and average da / dt (K = 33 ) (In the lobe with Ni content of 27% to 32%) is used to form the FIG. 8 graph.

Preferably K = 33 And crack growth rate (da / dt) at a temperature of 538 ° C. is 1 × 10 mm / sec or less. This is an approximate da / dt of INCOLOY alloy 909 annealed at 1800 ° F or 982 ° C in fine grain conditions. It is more advantageous if da / dt is less than 5 × 10 mm / sec in the dl state, and the approximate da / dt of INCONEL alloy 718 follows fine grain, delta-precipitation annealing (1750 ° F-1800 ° F, 954 ° C-982 ° C). . The reduction of crack growth is achieved in several ways through three different heat treatments, which have special advantages and disadvantages:

<1> Low Temperature Annealing (≤ 1850 ° F., 1010 ° C.): A crack growth rate of less than 10 × 10 ⁻⁵ inches / minute (4.2 × 10 mm / sec) is achieved at 1850 ° F. (1010 ° C.) and 5 × 10 ⁻ A growth rate of ⁵ inches / minute (2.1 × 10 ⁻⁵ mm / sec) is achieved upon annealing at 1800 ° F. (982 ° C.). Lower da / dt is possible with overaging (> 1450 ° F, 788 ° C) aging heat treatment.

(Advantages) 1) High yield strength at low annealing temperature:

2) da / dt is less sensitive to aging heat treatment and can select a wide range of aging temperatures.

3) Low temperature annealing is suitable for alloys such as alloy 718 (alloys connected to alloy 718 can be easily heat treated).

(Disadvantages) 1) It is sensitive to pre-thermomechanical processing.

2) Anisotropy of more coarse beta grains precipitated during treatment can lead to anisotropic mechanical properties.

3) More abundant and coarse beta particles cause ductile loss after prolonged medium temperature exposure.

4) Creep resistance is reduced due to fine grain and rich beta phase.

5) Not compatible with high temperature soldering heat treatment used to connect turbine engine casing and seal.

Low temperature annealing is performed for 0.5 to 10 hours. 0.5 to 6 hours are more preferred. The temperature should be at least 1650 ° F (900 ° C).

<2> high temperature beta aging treatment ≥ 1450 ° F. (788 ° C.): This aging temperature range is 10 × 10 ⁻⁵ inch / min (4.2 × 10 ⁻⁶ mm / sec), 5 ×, at all annealing temperatures. Effective at dropping to 10 ⁻⁵ (2.1 1 × 10 ⁻⁶ mm / sec) or less.

(Advantages) 1) If the aging temperature is ≧ 1500 ° F. (816 ° C.), it has good crack growth resistance regardless of the annealing temperature.

2) provides exceptional da / dt resistance for annealing at> 1850 ° F (1010 ° C) and <2000 ° F (1093 ° C).

3) Improve stress rupture ductility.

(Disadvantages) 1) Due to the richer beta phase and the larger grain boundary beta-matrix interface, it is prone to instability at 1000 ° F (538 ° C).

2) Creep Strength and Rupture Life Reduction (If Aging Time Is Not Short)

3) Heat treatment that is not always consistent with other material heat treatments in the combined engine parts.

Beta aging is appropriate for 0.5 to 24 hours, especially 1 to 6 hours. The beta aging treatment takes place at a temperature of 820 ° C or higher and 890 ° C or lower.

<3> High Temperature Annealing (> 2000 ° F., 1093 ° C.): The annealing at 2050 ° F. results in a da / dt ratio of 5 × 10 inches / minute (2.1 × 10 mm / second) or less.

(Advantages) 1) Many beta particles, including the primary beta phase, are dissolved and reprecipitation is controlled as fine particles at the grain boundaries.

2) The grain size is slightly granulated and the grain structure and isotropy of the remaining beta phase are restored.

3) The da / dt dependency on aging heat treatment temperature is reduced.

4) Excellent balance between stress rupture strength, creep resistance and da / dt resistance.

5) Provide optimal impact strength.

(Disadvantages) 1) Yield strength decrease, 2) Insufficient beta precipitation results in notched stress rupture fragments.

Hot annealing is carried out for 0.5 to 10 hours, especially for 0.5 to 6 hours and at temperatures below the melting point, in particular below 2125 ° F. (1163 ° C.).

The effect of heat treatment on room temperature tensile yield strength and elongation, and 649 ° C / 586 MPa combination smooth notch (Kt 3.7) stress rupture life and elongation is discussed.

Hot-rolled 30 is press-forged and lathed to 8 '(20 cm) diameter, hot rolled and hot rolled to form a 711 mm OD × 610 mm ID × 162 mm height gas turbine engine ring. Tension and stress rupture specimens are cut in the longitudinal direction. Flexible gauge bar tensile test is performed at 24 ° C in accordance with ASTM E8. Stress rupture tests were performed in medium to high humidity (30% to 60% relative humidity) using smooth notch (Kt 3.7) bars formed using standard low stress polishing techniques at 649 ° C. and a total cross-sectional stress of 586 MPa. Stress rupture tests and test specimens comply with ASTM E292.

Annealing at 1038 ° C and 1120 ° C produces a material with poor stress rupture life under gentle conditions. Water annealing after annealing produces very soft material and the material ages during slower air cooling. This age hardening is the result of beta and gamma-prime phase precipitation. However, if the furnace is cooled slowly to the settling temperature, sufficient strength is obtained, but such hardening does not result in sufficient tensile or stress rupture strength.

The study of the effect of heat treatment on 538 ° C da / dt shows that 1121 ° C annealing significantly improves da / dt resistance compared to 1038 ° C annealing (slower crack growth rate). When annealing at high temperatures, careful control is needed to prevent rapid grain growth that occurs above the beta dissolution temperature of 2070 ° F. (1130 ° C.). The annealing temperature between 1010 ° C. and 1090 ° C. does not dissolve a sufficient amount of beta and therefore limits Al available for new beta reprecipitation in a controlled manner when using other aging heat treatments. As a result, the mechanical properties of the annealing material at this temperature tend to change slightly with aging heat treatment and require longer aging heat treatment for longer cracking resistance (> 800 ° C and> 12 hours exposure).

The focus of this discussion on the mechanical properties is based on 1121 ° C. annealing. Hot annealing is the dissolution of a sufficient amount of beta and almost all gamma-frames, spherical and spherical beta particles remaining during the dissolution of the martensite phase, resulting in the dissolution of Al into the gamma matrix. Further dissolved Al is available for reprecipitation during the aging heat treatment according to the aging heat treatment used as intragranular fine spherical beta, discrete and fine intergranular beta or gamma-prime.

1121 ℃ Annealed + Isothermal Aging

After 1121 ° C. annealing, isothermal aging at 732 ° C. to 843 ° C. shows several results.

8 hours isothermal aging (732 ° C.) increases the yield strength by a useful strength level of 84 MPa to 644 MPa. However, stress rupture life and ductility decrease. At this temperature, the aging treatment is below the beta precipitation temperature, which precipitates gamma-prime but does not produce beta. In addition, the pre-treated precipitated spherical beta particles exhibit degradation close to the DO ₃ configuration, which is very similar to that found in Fe ₃ Al, and small amounts of flaps are formed in the spherical beta particles at the beta-gamma interface and the beta-betagrain interface. Although not confirmed positively, this flap is a martensitic BCT based on Ni ₅ Al ₃ or Ni ₂ Al.

Thus, with such heat treatment, the material is greatly improved in strength, but the stress rupture life and ductility are worse because they are susceptible to cracks accelerated by oxygen. The presence of the crescent intergranular rupture zone adjacent to the soft intergranular tensile rupture zone on the notched rupture surface shows rapid crack growth due to the grain boundary oxygen erosion accelerated by stress.

2. The 16-hour aging treatment at 788 ° C has excellent stress rupture life and ductility, yield strength increases (31 MPa) but is below the desired level. This temperature is above the minimum beta precipitation temperature or below the gamma-prime melting temperature as in FIG. The gamma prime phase should precipitate below the gamma prime melting temperature of 1500 ° F. (815 ° C.). The resulting microstructure includes both newly precipitated beta and gamma primes in addition to the beta spheres. However, the gamma-prime particles are relatively coarse due to higher settling temperatures and longer exposure times, so the yield strength increase is minimal. The combination of precipitated beta particles (which occur both inside the grain and at the boundaries) and coarse gamma-prime (which leads to greater micro-creep firing) results in very good burst life and ductility by inhibiting crack growth. However, yield strength is unsuitable for applications requiring high strength.

3. Aging at 843 ° C. results in a degraded but still acceptable stress rupture life and excellent ductility, but yield strength is reduced to lower levels than annealing and air cooled materials. This temperature is higher than the gamma-prime melting temperature (Figure 9) and is within the beta precipitation temperature range. Large amounts of beta precipitate out of the solid solution at the grain boundaries and within the grains due to the conversion of gamma prime to beta. Gamma-prime particles do not convert or dissolve into beta, so they are granulated in size and are not effective as tensile enhancers. The result is an acceptable 649 ° C. burst life and excellent ductility but lower yield strength than an annealed material (which shows good resistance to cracking), making it unsuitable for applications requiring high strength.

More preferably, 1 to 30 hours isothermal aging is performed after annealing treatment for 10 to 10 hours at the melting point of 1010 ° C to the alloy. In particular isothermal annealing between 1350 ° F and 1500 ° F (732 ° C and 815 ° C). Isothermal aging provides stress rupture strength and life with some ductile losses.

1121 ℃ Annealed + Two Stage Heat Treatment

Effects of furnace cooling, hold for 8 hours and air cooling to 621 ° C and 788 ° C aging heat treatment and 621 ° C at 56 ° C / h rate are discussed.

1.732 ° C./8h FC 621 ° C./8h AC.

Yield strength is significantly increased (105 Ma) compared to isothermal 732 ° C. aging due to gamma-prime precipitation hardening by transformation and degradation in pre-treated beta particles. By en step heat treatment, the gamma-prime in the sample has a bimodal size distribution that enhances the tensile strength. When using two-step gamma prime aging heat treatment, it is important to optimize the yield strength by performing furnace cooling between the aging stages. However, the furnace cooling rate between the aging stages cannot be confirmed to have a measurable effect.

Gamma prime precipitation occurs during the aging process between 950 ° F and 1500 ° F (510 ° C and 815 ° C). Granulated gamma prime is precipitated at an aging temperature of 1250 ° F. to 1450 ° F. (677 ° C. to 788 ° C.). The fine gamma prime phase precipitates predominantly at temperatures between 1000 ° F. and 1300 ° F. (538 ° C. and 704 ° C.). The first and second gamma prime aging steps are preferably carried out for 0.5 to 12 hours, in particular for 1 to 10 hours.

However, gamma-prime precipitation does not contribute to stress accelerated grain boundary oxygen embrittlement and pre-beta precipitation is inadequate and thus stress burst life is sensitive to notch rupture that is sensitive to the environment. This heat treatment is suitable for room temperature applications requiring high strength but is not useful in elevated temperature applications.

2. 788 ° C / 16h FC 621 ° C / 8h AC.

Yield strength is greatly increased (162MPa) and is almost similar to the two-stage heat treatment at 732 ℃. In contrast to the two-stage heat treatment at 732 ° C, stress rupture life and ductility are excellent, which means significantly improved oxygen brittleness resistance and excellent crack growth resistance. The material under these conditions shows gamma prime precipitation in bimodal size distribution at the grain boundaries, with a significant amount of beta precipitation inside the grains and a finer beta precipitation at the grain boundaries.

Combining the gamma-prime and beta phases of mixed size distributions achieves a combination of high strength and excellent stress rupture life and ductility. This is an advantageous heat treatment for both room temperature and elevated temperature applications, including gas turbine engines.

Annealing at 1121 ° C + 3-stage aging heat treatment.

This heat treatment combines conventional gamma-prime or gamma-two-prime aging heat treatment, such as the aging treatment often used for INCOLOY alloy 909 or INCONEL alloy X750 or 718 with high temperature beta-precipitation heat treatment (843 ° C./2 h AC). High strength, excellent stress rupture life and ductility were obtained. In fact, much higher yield strengths are achieved compared to two-stage aging heat treatment.

The microstructure of this material has assembled gamma grains (ASTM # 5 to # 1) containing hexagonal gamma-prime with bimodal size distribution. Inside the grains, beta spheres formed during pretreatment and newly precipitated beta particles are found. Coarse beta spheres and particles show platelet phases within the beta matrix interface and beta-meta grain boundaries, and DO ₃ phases similar to Fe ₃ Al (pre-precipitated beta spheres interconnected by grain boundaries are found).

A three-step heat treatment utilizing shorter, higher temperature beta precipitation heat treatment shortens the total aging heat treatment time from 27 hours of 621 ° C / 8h AC heat treatment to less than 20 hours at 788 ° C / 16h FC 55 ° C / h. Moreover, short-term beta-precipitation heat treatment gives flexibility to gamma-prime aging heat treatment, so that the alloy can be conveniently heat treated when bonded to dissimilar superalloys such as INCONEL alloy 706 or 718. Moreover, this alloy may be bonded or plated with a ceramic such as silicon nitride. Table 16 summarizes the mechanical test data resulting from the heat treatment.

TABLE 16

Notes: 1) AC = air cooled to room temperature

WQ = water cooled to room temperature

FC = furnace cooling 56 ° C / h up to the shown temperature

2) NT = not tested

3) YS = 0.2% offset yield strength, EL = elongation

4) Notch = notch rupture at the shown life time

TABLE 17

Add 0.007% to each slice.

As the titanium content changes, the effect of long-term exposure on the stability of the compositions in Table 17 is tested.

TABLE 18

TABLE 19

Basic heat treatment

1121 ° C./1 h. AC + 843 ℃ / 2h, AC + 718 ℃ / 8h FC (38 ℃ / h) to 621 ℃ / 8h, AC

In Table 18, the alloy obtains a small amount of strength without significant loss of ductility after 538 ° C exposure. The intensity is constant after 649 ° C. exposure and decreases slightly after 704 ° C. exposure. However, Alloy 6 of 5.5% Cr and 0.5% Ti shows brittleness after 1000 hours of exposure at 704 ° C. From the data it can be seen that it is advantageous to limit the titanium content to 0.5% by weight or less.

In FIG. 10, under these heat treatment conditions, da / dt of segment 10 was improved over alloy 909 and more than doubled over similar alloys without chromium. (49.5 Similar to 718 alloy at stresses below).

The tenth alloy was annealed at 1121 ° C. for 1 hour, air cooled, beta precipitate at 843 ° C. for 1 hour, air cooled, 1 hour at 732 ° C. for 2 hours gamma prime aging, furnace cooled to 641 ° C., 1 hour. Maintenance, and also air cooling. There is some directional effect on da / dt. However, the two curves are within the da / dt test precision range and are not very different. These data show how the annealing and aging heat treatment effects combine to achieve the desired useful properties.

The alloy of the present invention is suitable for most casting applications. Similar alloys show desirable casting properties. Beta phase formation also provides good weldability for high Al-containing alloys (high Al superalloys are difficult to weld). The alloy of the present invention can be formed by powder metallurgy mechanical alloys or thermal spray precipitation of oxide dispersions such as powder metallurgy, yttria.

As will be appreciated by those skilled in the art, various modifications may be made using the features herein.

Claims (9)

By weight ratio 26-50% cobalt, 20-40% nickel, 20-35% iron, 4-10% aluminum, 0.5-5% total niobium and 1/2 tantalum, 1.5-10% chromium, less than 1 titanium, Less than 0.2% carbon, less than 1% copper, less than 2% manganese, less than 2% silicon, less than 8% molybdenum, less than 8% tungsten, less than 0.3 boron, less than 2% hafnium, less than 2% Rhenium, less than 0.3% zirconium, less than 0.5% nitrogen, less than 1% yttrium, less than 1% rare earth metals except less than 1% lanthanum, less than 1% cerium, less than 1% magnesium, less than 1% Of calcium, oxide dispersion of less than 4% and other impurities, 538 ° C, 33 Alloy of controlled coefficient of thermal expansion having a crack growth rate of less than 1 × 10 ⁻⁴ mm / s at a stress intensity of. The cobalt content is 28-45%, the nickel content is 25-35%, the iron content is 22-30%, the aluminum content is 4-8%, the niobium and 1/2 tantalum are 1-4%, Alloy characterized by 1.5-5% chromium, less than 0.5% titanium and less than 0.1% carbon. The method according to claim 1, which has a body centered cubic beta phase obtained by annealing and an intermediate temperature aging treatment and a gamma prime phase obtained by aging treatment. Alloy characterized by having a static crack life of at least 10 hours at the initial stress intensity of. The yield strength of claim 690 MPa at least 0.2%, the elongation at least 10% at room temperature, the yield strength at least 590 MPa at 0.2% at 704 ° C., the elongation at least 15% at 704 ° C., the yield strength of at least 379 Ma at 649 ° C. An alloy characterized by having a Charpy V-notch impact energy of 5 N.M or more at room temperature and a coefficient of thermal expansion of 13.6 µm / m / ° C or less at 649 ° C. The cobalt content is 30-38%, the nickel content is 26-33%, the iron content is 24-28%, the aluminum content is 4.8-6.0, and the niobium and 1/2 tantalum are 2-3.5%. And chromium being 2-4%. 30-38% Cobalt, 26-33% Nickel, 24-28% Iron, 4.8-6.0% Aluminum, 2-3.5% Niomium + 1/2 Tantalum, 2-4% Chromium, Less Than 0.2% Titanium , Less than 0.5% carbon, less than 0.5% copper, 0.5% manganese, 0.5% silicon, less than 1% copper + manganese + silicon, less than 3% molybdenum, less than 3% tungsten, less than 5% molybdenum + tungsten , Less than 0.015% boron, less than 0.5% hafnium, less than 0.5% rhenium, less than 0.1% zirconium, less than 0.2% bilso, less than 0.2% yttrium, 0.2% lanthanum, less than 0.2% rare earth except lanthanum 538 ° C, 33 as an alloy consisting of metal, less than 0.2% cerium, less than 0.2% magnesium, less than 0.2% calcium, less than 0.2% oxide and other impurities Alloy of controlled thermal expansion coefficient having a crack growth rate of less than 1 × 10 ⁻⁴ mm / s at stress strength. Weight ratio 26-50% cobalt, 20-40% nickel, 20-35% iron, 4-10% aluminum, 0.5-5% total niobium and 1/2 tantalum, 1.5-10% chromium, less than 1% titanium , Less than 0.2% carbon, less than 1% copper, less than 2% manganese, less than 2% silicon, less than 8% molybdenum, less than 8% tungsten, less than 0.3% boron, less than 2% hafnium, 2 Less than% rhenium, less than 0.3% zirconium, less than 0.5% nitrogen, less than 1% yttrium, less than 1% lanthanum, less than 1% rare earth metals, less than 1% cerium, less than 1% magnesium , Annealing alloys of less than 1% calcium, less than 4% oxide dispersion and other impurities, annealing at temperatures between 1066 ° C. and below the melting point of the alloy or below 1010 ° C. and aging to temperatures below 815 ° C. for gamma prime precipitation. Alloy heat treatment method characterized in that. Weight ratio 26-50% cobalt, 20-40% nickel, 20-35% iron, 4-10% aluminum, 0.5-5% total niobium and 1/2 tantalum, 1.5-10% chromium, less than 1% titanium , Less than 0.2% carbon, less than 1% copper, less than 2% manganese, less than 2% silicon, less than 8% molybdenum, less than 8% tungsten, less than 0.3% boron, less than 2% hafnium, 2 <% Rhenium, <0.3% zirconium, <0.5% nitrogen, <1% yttrium, <1% lanthanum, <1% rare earth metals except lanthanum, <1% cerium, <1% magnesium Aging consisting of less than 1% calcium, less than 4% oxide dispersion and other impurities can be aged at temperatures between 788 ° C. and 890 ° C. for beta phase precipitation and at temperatures below 815 ° C. for gammaprime precipitation. Alloy heat treatment method characterized in that. Weight ratio 26-50% cobalt, 20-40% nickel, 20-35% iron, 4-10% aluminum, 0.5-5% total niobium and 1/2 tantalum, 1.5-10% chromium, less than 1% titanium , Less than 0.2% carbon, less than 1% copper, less than 2% manganese, less than 2% silicon, less than 8% molybdenum, less than 8% tungsten, less than 0.3% boron, less than 2% hafnium, 2 <% Rhenium, <0.3% zirconium, <0.5% nitrogen, <1% yttrium, <1% lanthanum, <1% rare earth metals except lanthanum, <1% cerium, <1% magnesium , An alloy consisting of less than 1% calcium, less than 4% oxide dispersion and other impurities are annealed at a temperature of 1010 ° C. to alloy melting point and 1-30 for beta and gamma prime particle precipitation at temperatures of 732 ° C. to 815 ° C. Alloy heat treatment method characterized by isothermal aging for a time.