Formation of an Intermediate Layer Between Grains in Nickel-Based Superalloy Turbine Blades
The boundary region formed on the surface of nickel-based single-crystal turbine blades was investigated by high-resolution microscopy observation. There was a distinguishable intermediate layer with the size of about 2 to 5 μm between the matrix and surface defect grains such as stray grains, multiple grains, freckle grains, and even low-angle grain boundaries which were formed during the solidification of turbine blades. The intermediate layer was composed of many elongated γ′ as well as γ phases. In addition, only one side of the intermediate layer was coherent to the matrix grain or defect grain due to good orientation match. At the coherent interface, the γ′ (as well as γ) phase started to extend from the parent grain and coincidently, rhenium-rich particles were detected. Furthermore, the particles existed within both elongated gamma prime and gamma phases, and even at their boundary. Based on experimental observations, the formation mechanism of this intermediate layer was discussed.
KeywordsTurbine Blade Intermediate Layer Solution Heat Treatment Discontinuous Precipitation Coherent Interface
As superalloys have excellent mechanical strength, phase and surface stability, and resistance to creep, corrosion, and oxidation under relatively severe mechanical stresses at elevated temperatures close to their melting point,[1, 2, 3, 4] they have found widespread application in gas turbine engines for jet propulsion and electricity generation. Newly developed superalloys are continually sought to be used in the hottest parts of the engine because the efficiency of the engine, fuel economy and reduction of emissions, can be improved by higher operation temperatures. The strengthening mechanism for these superalloys is mainly by solid solution hardening and the precipitation of an intermetallic phase. Solid solution hardening is achieved by the addition of different soluble elements, such as Cr, Mo, W, and Re via the inhibition of dislocation movement and the decrease of stacking fault energy in the crystal lattice, which leads to the inhibition of cross slip of dislocations. Precipitation hardening is obtained through the additions of Al, Ti, and Nb which have limited solubility in the alloy matrix. During heat treatment, a supersaturated solid solution generates finely distributed precipitates of gamma prime (γ′) phase which inhibits dislocation movement.
Superalloys are one of the most compositionally complex alloys developed, sometimes containing more than ten alloying elements, and through their component manufacture they are subjected to multiple steps of melting, casting, and heat treatment. As a result, they are subject to numerous metallurgical phenomena, such as melting, solidification, homogenization, aging, precipitation, transformation, coarsening, microsegregation, and chemistry variation. Therefore, these complexities can make it difficult to understand exactly the strengthening mechanism of Re in superalloys. The distribution of Re clusters, about 1 nm across, in the γ phase, which was observed by field ion microscopy and atom probe, may induce this strengthening mechanism,[5,6] while there are other studies showing no clustering of Re in superalloys, which was confirmed by extended X-ray absorption fine structure (EXAFS), atom probing, and first-principles density functional theory calculations.[7,8] Depending on the analysis method or equipment, different interpretations can be made. In addition, as analysis techniques and equipments develop, new findings can be made even in these widely investigated alloy systems. One of these recent findings is the formation of an intermediate layer containing Re-rich particles in Ni-based single-crystal turbine blades along grain boundaries. In single-crystal superalloys, the reintroduction of any grain boundaries can dramatically reduce the superior mechanical properties because grain boundary strengthening elements, such as boron, carbon, and zirconium, have been removed[3,4] and more importantly, creep rupture can occur along the boundary. In addition, if Re-rich particles form, they lose their main role as a solid solution hardening element which distorts the atomic lattice of the gamma (γ) phase matrix and inhibits dislocation movement. Therefore, the existence of any boundary and this associated intermediate layer, and the formation of Re-rich particles in a single-crystal superalloy are undesirable. In this study, based on this previous reporting of the intermediate layer in turbine blades and the detection of Re-rich particles in a sample containing a stray grain, the boundary regions between matrix grains and several other defect grains, such as stray grains, multiple grains, freckle grains, and even low-angle grain boundaries found in nickel-based single-crystal turbine blades have been investigated for better understanding of these materials. In addition, high-resolution analysis has been carried out from several other Ni-based single-crystal alloys, such as SRR99 and CMSX-4, as well as CMSX-10. Based on experimental observation, the formation mechanism of the intermediate layer, and the relationship between Re-rich particles and the layer have been also discussed.
Figures 2 and 3 also show that the size of the boundary layer is about 2 to 5 μm, which means that it is a layer (not an independent grain boundary) formed between adjacent two grains because the size of effective grain boundaries in other systems is shown to be about 0.5 nm. Most of all, the observation suggests that the intermediate layer can be detected in any grain boundary regions formed during the solidification of Ni-based single-crystal turbine blades even though the boundary width is slightly different. In addition, close observation of the intermediate layer formed between a stray grain and the matrix in the platform region (Figures 2 and 3(a)) shows that the stray grain is surrounded with the intermediate layer, and as a result, the boundary was easily distinguishable during observation. It is interesting to note that only one side of the layer marked with an arrow in Figure 3(a) has a coherent boundary. Most of all, the coherent interface could be found on both the sides of the stray grain or the side of the matrix separately, which means that there is no tendency to form the coherent interface in any particular grain. The observation suggests that one side of the intermediate layer has a coherent interface with one of the grains.
4.1 Possibility of Cellular Recrystallization or Formation of Columnar Zone of Elongated Grains
The elongated γ′ phase in the intermediate layer is similar to the lamellar microstructure formed by cellular recrystallization,[26, 27, 28, 29, 30] or the columnar zone of elongated grains formed in solidified alloy ingots. The samples in this study were subjected to multiple steps of melting, casting, and heat treatment during manufacturing. The solution heat treatment process undertaken on these samples was intended to dissolve the γ/γ′ eutectics and reduce the chemical segregation resulting from the dendritic growth during the solidification process. Therefore, even though the elongated γ′ phase might form during solidification, it would have been dissolved as the heat treatment was above the γ′ solvus, and as a result there was little possibility that the elongated γ′ phase was analogous to the columnar zone of elongated grains formed during solidification.
During the manufacturing and processing of new parts, plastic deformation may be induced in the component from many sources, such as contracting stresses during cooling of the solid metal in the shell mold, mechanically removing shell mold residue, removal of the part from the gating system and general handling, and these can be concentrated on the surface of components or even extend into the bulk of the part. This concentrated deformation could be released by recrystallization during solution heat treatment and subsequent grain growth by strain-induced boundary migration.[31,32] As a result, cellular recrystallization which is similar to the elongated γ′ and γ phases has been observed in many specimens.[26, 27, 28, 29, 30, 31, 32] However, in order to initiate the recrystallization, a free surface as well as γ′-free zone must be available and this allows the recrystallization to occur on the surface and move into the bulk of the material. If these conditions do not exist, then a much higher amount of plastic deformation is required (e.g., more than about 10 pct). After that, the cellular recrystallization moves a grain boundary into the deformed parent material. The reduction of free energy (ΔG T) by the cellular recrystallization can be expressed as ΔG T = ΔG m + ΔG γ + ΔG c, where ΔG m reflects the annihilation of dislocations, ΔG γ is related to the microstructural coarsening and creation of a high-angle grain boundary, and ΔG c is the chemical composition of phases in the recrystallized region. As our previous study showed that the compositions of γ′ in the elongated region and the matrix or a stray grain are almost same, the last term is negligible. The realistic strain values which can be imposed on the components are about 2 to 3 pct. Therefore, the value of ΔG T is positive without any additional deformation, such as a hardness indentation or similar non-uniform deformation in excess of about 10 pct, which means that it is almost impossible to induce the recrystallization in the bulk materials without additional deformation. In the previous study an investigation into a recrystallized grain showed that the elongated γ and γ′ were not present despite the boundary moving through the material to the final location during solution heat treatment. In addition, it should be emphasized that in the sample observed in this study, the top surface of the component (which plastic deformation was concentrated on) was etched away for visual inspection, and the elongated direction of the intermediate layer was nearly parallel to the surface through the thickness of the component. Therefore, the elongated γ′ phase in the intermediate was not formed either by the cellular recrystallization or by the formation of a columnar zone of elongated grains during solidification.
4.2 Formation Mechanism of Intermediate Layer
Before discussing the formation mechanism of the intermediate layer, it is necessary to clarify whether the sharp (non-coherent) interface is the starting point in the formation of the layer or the terminating one. To form the elongated phases, either γ′ or γ should be nucleated. For nucleation, the activation energy should be reduced through releasing some free energy (ΔG T) which can be expressed as ΔG T = Aγ−V(ΔG v−ΔG s), where A and V are area and volume, respectively, γ an interfacial energy, ΔG v a volume free energy, and ΔG s is a misfit strain energy per unit volume. Therefore, to achieve a low interfacial energy, it is important to achieve a good lattice match through a small disregistry (δ = Δa 0/a o), where Δa 0 is the difference between the lattice parameters of the nucleating and the nucleated metal belonging to the same crystal structure along a specific direction, and having similar crystallography.[33, 34, 35] At this point, it should be emphasized that the sharp interface between the elongated phase and the matrix grain (or a stray grain) means that there was no specific orientation relationship between them. Therefore, the non-coherent interface was not energetically the starting point of the formation of the intermediate layer, which means that the formation of the layer started from the coherent interface and terminated at the topologically sharp one.
4.3 Comparison of the Formation Mechanism with a Conventional Discontinuous Precipitation Reaction
4.4 Relationship Between the Intermediate Layer and Re-rich Particles
5 Summary and Conclusion
An intermediate layer with a width of about 2 to 5 μm containing fine Re-rich particles was detected between a matrix grain and several surface defect grains which formed during solidification. The intermediate layer was composed of many elongated γ′ as well as γ phases (not purely γ′ phase) and had two different interfaces: coherent and non-coherent sharp interfaces. From the side of the coherent interface, γ′ (as well as γ) phase initiated and grew in elongated form, not the cuboidal structure seen in the body of the grain. Most of all, the elongated γ′ phase formed from γ′ precipitates in an adjacent surface defect grain or matrix and retained the orientation of these initiating grains. This morphology of the intermediate layer was similar to an interfacial layer formed from the discontinuous precipitation. However, in the intermediate layer, some Re-rich particles existed along only one side which was coherent and they existed in both the γ′ phase and the γ phase, and even at their boundary. Most of all, the Re-rich particles existed from as-cast samples. Re-free or samples containing medium amount of Re clearly supported that the elongation of γ′ and γ phases in the intermediate layer was definitely related to the existence of Re-rich particles.
It is well known that grain boundaries are a source of weakness of turbine blades.[3,4] The existence of grain boundaries, exactly, the existence of the intermediate layer is probably detrimental to the superior mechanical properties of Ni-based turbine blades. Therefore, even though any direct measurements of mechanical properties was not done in this study, it is suggested that it is necessary to investigate the methods to minimize or avoid the formation of the intermediate layer as well as numerous Re-rich particles.
The financial support and provision of evaluation test pieces by Rolls-Royce is acknowledged.
- 1.M. Durand-Charre, The Microstructure of Superalloys (OPA, Amsterdam, 1997), pp. 1–14Google Scholar
- 3.B. Geddes, H. Leon, X. Huang, Superalloys: Alloying and Performance (ASM International, Materials Park, 2010), pp. 1–15Google Scholar
- 10.W.S. Walston, J.C. Schaeffer, W.H. Murphy, Superalloys 1996 (Minerals, Metals & Materials Soc, Warrendale, 1996), pp. 9–18Google Scholar
- 11.G.L. Erickson, Superalloys 1996 (Minerals, Metals & Materials Soc, Warrendale, 1996), pp. 35–44Google Scholar
- 16.Energy table for EDS analysis (JEOL, 2013), http://www.jeol.co.jp. Accessed 28 Sept 2013.
- 25.D.A. Porter, K.E. Eastering, M.Y. Sherif, Phase Transformations in Metals and Alloys, 2nd edn. (CRC Press, Boca Raton, 2009), pp. 322–25Google Scholar
- 46.W.S. Walston, K.S. O’Hara, E.W. Ross, T.M. Pollock, W.H. Murphy, Superalloys 1996 (Minerals, Metals & Materials Soc, Warrendale, 1996), pp. 27–34Google Scholar
Open AccessThis article is distributed under the terms of the Creative Commons Attribution 4.0 International License (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use, distribution, and reproduction in any medium, provided you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons license, and indicate if changes were made.