Patent Application
20060127695
This invention relates to methods for depositing alloy compositions for high-temperature, oxidation resistant coatings. Coatings based on these alloy compositions may be used alone or, for example, as part of a thermal barrier system for components in high-temperature systems.
The components of high-temperature mechanical systems, such as, for example, gas-turbine engines, must operate in severe environments. For example, the high-pressure turbine blades and vanes exposed to hot gases in commercial aeronautical engines typically experience metal surface temperatures of about 900-1000° C., with short-term peaks as high as 1150° C. A portion of a typical metallic article 10 used in a high-temperature mechanical system is shown in
The adhesion and mechanical integrity of the TGO scale layer 18 is very dependent on the composition and structure of the bond coat 16. Ideally, when exposed to high temperatures, the bond coat 16 should oxidize to form a slow-growing, non-porous TGO scale that adheres well to the superalloy substrate 12. Conventional bond coats 16 are typically either (i) an MCrAlY overlay (where M=Ni, Co, NiCo, or Fe) having a β-NiAl+γ-Ni phase constitution or (ii) a platinum-modified diffusion aluminide having a β-NiAl phase constitution. The Al content in either of these types of coatings is sufficiently high that the Al2O3 scale layer 18 can “re-heal” following repeated spalling during service of the turbine component.
However, as a result of this Al enriched composition and the predominance of the β-NiAl in the coating microstructure, these coatings are not compatible with the phase constitution of the Ni-based superalloy substrates, which have a gamma-Ni phase and a gamma prime-Ni3Al (referred to herein as γ-Ni+γ′-Ni3Al or γ+γ′) microstructure. When applied to a superalloy substrate having a γ-Ni+γ′-Ni3Al microstructure, Al diffuses from the coating layer to the substrate. This Al interdiffusion depletes Al in the coating layer, which reduces the ability of the coating to sustain Al2O3 scale growth. Additional diffusion also introduces unwanted phase changes and elements that can promote oxide scale spallation. A further drawback of β-NiAl-based coatings is incompatibility with the γ-Ni+γ′-Ni3Al-based substrate due to CTE differences.
Another approach to depositing a protective coating on a γ-Ni+γ′-Ni3Al-based metallic article 28, described in U.S. Pat. Nos. 5,667,663 and 5,981,091 to Rickerby et al., is shown in
Copending application U.S. Ser. No. 10/439,649, incorporated herein by reference, describes alloy compositions suitable for bond coat applications. The alloys include a Pt-group metal, Ni and Al in relative concentration to provide a γ+γ′ phase constitution, with γ referring to the solid-solution Ni phase and γ′ referring to the solid-solution Ni3Al phase. In these alloys, a Pt-group metal, Ni and Al, are present, and the concentration of Al is limited with respect to the concentrations of Ni and the Pt-group metal such that the alloy includes substantially no β-NiAl phase. These alloys are shown in the region A in
Preferably, the ternary Ni—Al—Pt alloy in the copending '649 application includes less than about 23 at % Al, about 10 at % to about 30 at % of a Pt-group metal, preferably Pt, and the remainder Ni. Additional reactive elements such as Hf, Y, La, Ce and Zr, or combinations thereof, may optionally be added to or present in the ternary Pt-group metal modified γ-Ni+γ′-Ni3Al alloy and/or improve its properties. The addition of such reactive elements tends to stabilize the γ′ phase. Therefore, if sufficient reactive metal is added to the composition, the resulting phase constitution may be predominately γ′ or solely γ′. The Pt-group metal modified γ-Ni+γ′-Ni3Al alloy exhibits excellent solubility for reactive elements compared to conventional β-NiAl-based alloys, and in the '649 application the reactive elements may be added to the γ+γ′ alloy at a concentration of up to about 2 at % (˜4 wt %). A preferred reactive element is Hf. In addition, other typical superalloy substrate constituents such as, for example, Cr, Co, Mo, Ta, and Re, and combinations thereof, may optionally be added to or present in the Pt-group metal modified γ-Ni+γ′-Ni3Al alloy in any concentration to the extent that a γ+γ′ phase constitution predominates.
The Pt-group metal modified alloys have a γ-Ni+γ′-Ni3Al phase constitution that is both chemically, physically and mechanically compatible with the γ+γ′ microstructure of a typical Ni-based superalloy substrate. Protective coatings formulated from these alloys will have coefficients of thermal expansion (CTE) that are more compatible with the CTEs of Ni-based superalloys than the CTEs of β-NiAl-based coatings. The former provides enhanced coating stability during the repeated and severe thermal cycles experienced by mechanical components in high-temperature mechanical systems.
When thermally oxidized, the Pt-group metal modified γ-Ni+γ′-Ni3Al alloy coatings grow an α-Al2O3 scale layer at a rate comparable to or slower than the thermally grown scale layers produced by conventional β-NiAl—Pt bond coat systems, and this provides excellent oxidation resistance for γ-Ni+γ′-Ni3Al alloy compositions. When the Pt-metal modified γ+γ′ alloys further modified with a reactive element such as, for example, Hf, and applied on a superalloy substrate as a coating, the growth of the TGO scale layer is even slower than comparable coating compositions without Hf addition. After prolonged thermal exposure, the TGO scale layer further appears more planar and has enhanced adhesion on the coating layer compared to scale layers formed from conventional β-NiAl—Pt coatings.
In addition, the thermodynamic activity of Al in the Pt-group metal modified γ-Ni+γ′-Ni3Al alloys can, with sufficient Pt content, decrease to a level below that of the Al in Ni-based superalloy substrates. When such a Pt-group metal modified γ-Ni+γ′-Ni3Al alloy coating is applied on a superalloy substrate, this variation in thermodynamic activity causes Al to diffuse up its concentration gradient from the superalloy substrate into the coating. Such “uphill diffusion” reduces and/or substantially eliminates Al depletion from the coating. This reduces spallation in the scale layer, increases the long-term stability of the coating and scale layers, and would greatly enhance the reliability and durability of a thermal barrier coating system.
The Pt-group metal modified γ-Ni+γ′-Ni3Al alloy may be applied to a superalloy substrate using any known process, including for example, plasma spraying, chemical vapor deposition (CVD), physical vapor deposition (PVD) and sputtering to create a coating and form a temperature-resistant article. Typically this deposition step is performed under non- or minimal oxidizing conditions.
As described earlier, when the Pt-group metal modified γ+γ′ alloys described in the '649 application are formulated with other reactive elements such as, for example, Hf, and applied on a superalloy substrate as a coating, the growth of the TGO scale layer is even slower than comparable coating compositions without Hf addition. After prolonged thermal exposure, the TGO scale layer further appears more planar and has enhanced adhesion on the coating layer compared to scale layers formed from conventional β-NiAl—Pt bond coat materials. As such, inclusion of a reactive element in the Pt-metal modified γ+γ′ alloys described in the '649 application is highly desirable.
As noted above, Rickerby et al. suggest that the reactive element Hf may be added to a Pt-metal modified γ+γ′ alloys at a level of up to 0.8 wt %, but providing a surface layer with a desired reactive element concentration has proved difficult. The reason for this is that the nearly complete partitioning of a reactive element such as Hf to the γ′ phase necessitates that γ′ be the principal phase during the deposition process to enrich the surface with Hf.
In one aspect, the invention is a method for making an oxidation resistant article, including (a) depositing a layer of a Pt group metal on a substrate to form a platinized substrate; and (b) depositing on the platinized substrate a reactive element selected from the group consisting of Hf, Y, La, Ce and Zr and combinations thereof to form a surface modified region thereon, wherein the surface-modified region comprises the Pt-group metal, Ni, Al and the reactive element in relative concentration to provide a γ-Ni+γ-Ni3Al phase constitution.
In preferred embodiments of this method, the surface modified region comprises greater than 0.8 wt % and less than 5 wt % of the reactive element. A preferred reactive element is Hf.
In another aspect, the invention is a method of making a temperature resistant article, including (a) depositing a layer of Pt on a superalloy substrate to form a platinized substrate; (b) heat treating the platinized substrate; and (c) depositing from a pack onto the platinized substrate to form a surface modified region thereon, wherein the pack comprises sufficient Hf such that the surface modified region includes Pt, Ni, Hf and Al in relative concentration to provide a γ-Ni+γ-Ni3Al phase constitution, and wherein the surface-modified region includes greater than 0.8 wt % and less than 5 wt % Hf.
In yet another aspect, the invention is a heat resistant article including a superalloy with a surface region including a reactive element selected from the group consisting of Hf, Y, La, Ce and Zr and combinations thereof, wherein the surface region includes a Pt-group metal, Ni, Al and the reactive element in relative concentration to provide a γ-Ni+γ′-Ni3Al phase constitution.
The Pt+reactive element-modified γ-Ni+γ′-Ni3Al coatings described herein have a number advantages over conventional β-NiAl containing coatings, including: (1) compatibility with the Ni-based superalloy substrate in terms of phase constitution and thermal expansion behavior; (2) no performance limiting phase transformations in the coating layer (i.e., destabilization of β to martensite or γ′) or in the coating/substrate interdiffusion zone (i.e., formation of brittle topologically close-packed (TCP) phases such as sigma); (3) the existence of a chemical driving force for the Al to diffuse up its concentration gradient from the substrate to the coating; (4) and exceptionally low TGO scale growth kinetics due, in part, to the presence of a preferred reactive element content of 0.8-5 wt %. Stemming from these advantages is the further advantage that the Pt+reactive metal-modified γ-Ni+γ′-Ni3Al coatings do not have to be as thick as the conventional β-NiAl containing coatings to provide a performance advantage.
The details of one or more embodiments of the invention are set forth in the accompanying drawings and the description below. Other features, objects, and advantages of the invention will be apparent from the description and drawings, and from the claims.
Like reference symbols in the various drawings indicate like elements.
In one aspect, the invention is a method for making an oxidation resistant article that includes an oxidation resistant region on a substrate, typically a superalloy substrate. The oxidation resistant alloy layer includes a modified γ-Ni+γ′-Ni3Al alloy containing a Pt-group metal, Ni, Al and a reactive element in relative concentration such that a γ-Ni+γ′-Ni3Al phase constitution results; although, stabilization effects by certain elements may cause γ′-Ni3Al to be the sole phase. In this alloy the concentration of Al is limited with respect to the concentration of Ni, the Pt-group metal and the reactive element such that substantially no β-NiAl phase, preferably no β-NiAl phase, is present in the alloy, and the γ-Ni+γ′-Ni3Al phase structure predominates.
The reactive element(s) in the oxidation resistant region tend not to oxidize even though their oxides are more stable than Al2O3. While not wishing to be bound by any theory, this is apparently because Pt acts to decrease the thermodynamic activity of Hf and Zr in the γ-Ni+γ′-Ni3Al. The oxidation resistant region may be formed on the substrate surface to impart oxidation and high-temperature degradation resistance to the substrate.
Referring to
Referring again to
As the layer of Pt-group metal 104 on the superalloy substrate 102 is heated, elements diffuse from the substrate 102 into the Pt-group metal region 104. This diffusion can continue until a γ-Ni+γ′-Ni3Al microstructure predominates within the Pt-metal group region 104. Thus, a diffusion heat-treatment preferably follows deposition of the Pt layer. As an example, the heat treatment may be for 1-3 hours at 1000-1200° C. During this heat treatment step, further diffusion occurs from the superalloy substrate 102 into the layer of Pt-group metal 104 to form a Pt-modified surface region in which γ′ is the principal phase, most preferably the sole phase. Current experimental data indicates that reactive elements such as Hf, Zr and the like partition almost solely to the γ′ phase. As a consequence, the full oxidative benefit of reactive element addition is most readily and easily realized when γ′ is the principal phase in the γ-Ni+γ′-Ni3Al microstructure of the region 104.
Referring to
In the pack-cementation process, for example, the substrate 102, including the Pt-group metal layer 104, are embedded in a powder mixture containing either a pure or alloyed coating-source material called the master alloy, a halide salt that acts as an activator, and a filler material.
During the deposition process, the powders in the pack are heated to an elevated deposition temperature, which produces a halide gas containing the reactive metal. When the Pt-group metal layer 104 is exposed to the reactive-metal-containing gas, the gas reacts with the layer 104, and the reactive metal deposits on the layer 104 to form a diffusion coating referred to herein as the surface modified region 106.
The composition of the surface modified region 106 is directly dependent on the composition of the powders in the pack. The pack powder composition preferably includes a filler, an activator and a master alloy source, and many compositions are possible. However, the pack powder composition should contain a sufficient amount of the master alloy source such that the reactive metal deposits on the Pt-group metal layer 104 and forms a surface-modified region 106 having the desired concentration of reactive metal. Preferably, the surface modified region 106 includes an average of up to about 5 wt % reactive metal, preferably about 0.8 wt % to about 5 wt %, and most preferably about 0.8% to about 3 wt %.
To achieve these concentrations of reactive metal in the region 106, typically the master alloy source includes at least about 1 wt % of a reactive metal, preferably Hf, and is present in the pack at a content of about 1 wt % to about 5 wt % Hf, but most preferably about 3 wt % Hf. A salt containing one or more of reactive-elements may be an alternative source, such as, for example, hafnium chloride. The master alloy source may optionally include about 0.5 wt % to about 1 wt % Al to provide surface enrichment of the Pt-metal layer 104.
The pack powder composition also includes about 0.5 wt % to about 4 wt %, preferably about 1 wt %, of a halide salt activator. The halide salt may vary widely, but ammonium halides such as ammonium chloride and ammonium fluoride are preferred.
The balance of the pack powder composition, typically about 94 wt %, is a filler that prevents the pack from sintering and to suspend the substrate during the deposition procedure. The filler typically is a minimally reactive oxide powder. Again, the oxide powder may vary widely, but compounds such as aluminum oxide, silicon oxide, yttrium oxide and zirconium oxide are preferred, and aluminum oxide (Al2O3) is particularly preferred to provide additional Al surface enrichment to Pt-metal layer 104.
The pack powder composition is heated to a temperature of about 650° C. to about 1100° C., preferably less than about 800° C., and most preferably about 750° C., for a time sufficient to produce a surface-modified region 106 with the desired thickness and reactive metal concentration gradient. The deposition time typically is about 0.5 hours to about 5 hours, preferably about 1 hour.
As the reactive metal and any other metals in the pack composition are deposited on the Pt-metal layer 104, diffusive mixing occurs at the surface of the layer 104 to form the surface modified region 106. The reactive metal, preferably Hf, as well as any other metals in the pack, such as Al, diffuse into and mix to form an Al-enriched Pt+reactive-metal modified γ-Ni+γ′-Ni3Al surface region 106. This surface-modified region 106 is therefore enriched in the metals from the pack. Within the surface-modified region 106, the concentration of reactive metal is greatest at the surface 107, and gradually decreases over the thickness of the layer 106, thus forming a reactive metal concentration gradient across the thickness of the layer 106.
The surface-modified region 106 typically has a thickness of about 5 μm to about 50 μm, preferably about 20 μm. Over the first 20 μm, the surface-modified region 106 has a composition including at least about 1 wt % of the reactive metal, preferably Hf, typically about 1 wt % Hf to about 3 wt % Hf.
During and after the deposition process, in addition to the inward diffusion from the surface modified region 106 into the Pt-group metal layer 104, metals also diffuse outward from the superalloy substrate 102 into the Pt-group metal layer 104 and further into the surface modified region 106. For example, a superalloy substrate 102 such as CMSX-4 nominally contains at least about 12 at % Al. The Al in the substrate diffuses into the Pt-group metal layer 104 and into the surface modified region 106. In addition, other elements from the superalloy substrate, such as, for example, Cr, Co, Mn, Ta, and Re may diffuse outward from the superalloy substrate 102 into the Pt-group metal layer 104 and then into the surface modified region 106. Further, if other metals such as Al are included in the pack, Al deposited along with the reactive metal layer may diffuse inward into the surface modified region 106 and into the Pt-group metal layer 104.
The composition of the pack is selected considering these outward and inward diffusive mixing behaviors, and it is important that while a variety of metals may be present in the surface modified region 106, the Al content of the region 106 is preferably controlled with respect to concentration of the Pt-group metal, Ni, and reactive element such that a γ-Ni+γ′-Ni3Al phase constitution results, with γ′-Ni3Al being the principal or even sole phase. In the region 106 the concentration of Al is limited with respect to the concentration of Ni, the Pt-group metal and the reactive element such that substantially no β-NiAl phase structure, preferably no β-NiAl phase structure, is present in the region, and the γ-Ni+γ′-Ni3Al phase structure predominates.
As a result of this extensive diffusive mixing, the amount of metallic Al as the master alloy source in the pack composition is preferably maintained at a very low level, less than about 1 wt %. Even master alloy sources including 0 wt % Al have been found to produce a γ-Ni+γ′-Ni3Al phase, particularly if the filler material includes at least some Al2O3 powder. The main source for Al in the surface modified region 106 can be the superalloy substrate 102, not the pack. Specifically, the chemical interaction between Al and Pt is such that a strong driving force exists for the Al to diffuse from the substrate 102 into Pt-group metal layer 104 and further into the surface modified region 106. Pack compositions with metallic Al concentrations of greater than about 1 wt % typically result in β-NiAl phase formation in the surface modified region 106, and often result in the formation of W-rich TCP precipitates therein.
The thickness of the Pt-group metal layer 104 also has an impact on the diffusive mixing behavior in the article 100, as well as on the composition of the surface modified region 106. For example, if the Pt-group metal layer 104 has a thickness of about 2 μm, the surface modified layer 106 most likely will have a Pt-group metal modified γ+γ′ coating with a primary γ phase, while a Pt-group metal layer with a thickness greater than about 4 μm, typically about 4 μm to about 8 μm, will most likely have a Pt-group metal modified γ+γ′ coating with a primary γ′ phase.
The temperature used in the pack cementation process also has an impact on the phase constitution of the surface modified layer 106. At higher temperatures, particularly when Al powder is included in the master alloy source, the amount of Al deposited along with the reactive metal becomes sufficiently high to produce unwanted β-NiAl phase structure in the surface modified region 106. Typically, a pack cementation temperature of about 900° C. resulted in some β-NiAl phase formation. Therefore, to reduce formation of β-NiAl phase structure in the surface modified region 106, the pack cementation temperature should preferably be maintained at less than about 800° C., preferably about 750° C.
Following the deposition process, the article 100 is preferably cooled to room temperature, although this cooling step is not required.
Following formation of the surface-modified region 106, the article 100 may optionally be heat treated at a temperature of about 900° C. to about 1200° C. for up to about 6 hours to stabilize the microstructure of the surface modified layer 200. The optional heat treatment step may be conducted prior to or before the article 100 is cooled to room temperature.
Referring to
Preferred embodiments of the invention will now be described with reference to the following examples.
An electrodeposition bath was prepared using a tetra-amineplatinum hydrogen phosphate solution ([Pt(NH3)4]HPO4). The superalloy substrate was CMSX-4 with approximate dimensions 15×10×1 mm.
The superalloy substrate sample was prepared by grinding to a 600-grit finish using SiC paper, followed by cleaning using the following procedure. First the sample was dipped in distilled water and dried with a tissue. The sample was then dipped in a 10 wt. % HCl solution for 2 minutes, dipped in distilled water and dried with a tissue. Finally, the sample was ultrasonically cleaned in acetone for 5 minutes and dipped in distilled water.
The prepared sample was then electrodeposited immediately. The electrodeposition conditions were as follows:
Current density ≈0.5 A/dm2
Temperature ≈95° C.
pH ≈10.5 (adjusted using NaOH)
Deposition time=0.5 hour
Distance between anode and cathode ≈5 cm
Anode: Pt
Anode:cathode surface area ratio ≈2
To produce a Pt+Hf-modified γ-Ni+γ′-Ni3Al coating in which γ′ was the principal phase, packs consisting of Hf powder and with/without Al powder were assessed. The basis for using no Al powder in the pack is that Al from the superalloy substrate will be driven to diffuse outward to the Pt-enriched surface, since Pt decreases the chemical activity of Al in γ and γ′ phase structures.
Using a pack deposition temperature of 750 or 800° C. and an NH4Cl activator content of about 1 wt %, it was found that Pt+Hf-modified coatings can be obtained. The following section will discuss the effects of specific experimental parameters on the microstructure and composition of Pt+Hf-modified coatings.
Thickness of Electrodeposited Pt Layer
By heat treating the Pt-coated sample, a simple Pt-modified coating can be obtained via inward Pt and outward Al+Ni diffusion. It was found that the thickness of deposited Pt layer affects the coating microstructure, composition and relative proportions of γ and γ′.
Al Content in Pack
The amount of aluminum powder in the pack will affect the extent of aluminum intake into the substrate. CMSX-4 nominally contains about 12 at % Al, which could also diffuse outward to the Pt-enriched surface during heat-treatment. Thus, it was deemed that only small amount of Al is required to obtain coating with about 22 at % Al by the pack cementation process.
It was also found that a Pt-coated CMSX-4 substrate that is further treated in a pack free of Al powder, yet still containing Al2O3 powder, can form a Pt-modified γ′-based surface layer.
Hf Content in the Pack
It is known that Hf partitions to the γ′ phase, and there must ultimately exist a critical Hf content in the pack to obtain a sufficiently high Hf deposition rate. From this study, it was found that 5 wt % Hf in the pack resulted in a detectable Hf content (above about 0.3 at %) in the γ+γ′ coating (see
Temperature of Pack Cementation Process
Temperature is a factor in determining the extent of Al deposition. At higher temperatures and using ˜1 wt % Al in the pack, the supply of Al becomes sufficiently high for unwanted (from the standpoint of obtaining a γ+γ′ coating) β-NiAl formation.
An aluminizing temperature above 900° C. resulted in dense β-NiAl formation, which was hard to transform to γ′ phase with heat-treatment, such as 1-4 days heat-treatment at 1100° C.
Referring to
A number of embodiments of the invention have been described. Nevertheless, it will be understood that various modifications may be made without departing from the spirit and scope of the invention. Accordingly, other embodiments are within the scope of the following claims.
