The oxidation kinetics of alumina-forming metals can be affected by adding a small amount of a reactive (normally rare earth) element oxide (RExOy) and the segregation of the reactive element (RE) ions to the growing alumina grain boundaries (GBs) has been considered as a responsible reason. However, this interpretation remains a controversial issue as to how RE ions are produced by RExOy which is thermodynamically and chemically stable in metals. The question is answered by a model that is based on transmission electron microscopy (TEM) investigation of a CeO2-dispersed nickel aluminide oxidized in air at 1100 °C. The CeO2 dispersion is incorporated into the alumina scale by the inward growth of inner α-Al2O3, where it partially dissolves producing tetravalent Ce cations which then transform to trivalent cations by trapping electrons. The trivalent cations segregate to the α-Al2O3 GBs and diffuse outward along first the GBs and later the twin boundaries (TBs) in the outer γ-Al2O3 layer, being precipitated as Ce2O3 particles near surface.
High temperature oxidation, a thermally- and chemically-activated reaction process with an expected increase in severity as the temperature increases, is a key mode of environmental degradation of high temperature structural materials. It normally leads to a loss of their service life. The service life can, however, be highly prolonged if the materials have the ability to develop a scale of thermally grown oxide (TGO) with the merits of compactness, slow growth and thermodynamic stability. The α-Al2O3 TGO is such a representative oxide, which can offer excellent resistance to oxidation above 1000 °C. Development and study of oxidation-resistant alumina formers at high temperatures has been attracting great interests for decades. Alumina-forming MCrAl system (M = Ni, Fe, Co, or their combinations) and nickel aluminides are the well-known oxidation-resistant alloys and coatings. Many reports show that these alumina formers can be further improved in the high temperature oxidation performance by adding small amounts of reactive elements such as yttrium, hafnium, cerium, zirconium and lanthanum1,2,3,4,5,6,7,8,9,10,11,12,13,14,15,16,17,18,19,20,21,22,23,24,25,26,27,28,29,30,31,32,33,34,35,36,37,38,39.
The phenomenon that the REs additions improve the oxidation resistance of metals was first found in 1937 by Pfeil40 and is popularly referred to as the “reactive element effects (REEs)”. The RE additions into alumina formers are conventionally made by alloying1,4,5,7,8,12,13,15,19,22,23,25,26,27,34,35,39, ion implantation5,8,9,11,15,29,31,35 and RExOy dispersions2,3,6,7,10,14,15,17,18,20,24,28,32,33. The REEs on alumina formers have been summarized and reviewed successively by some authors41,42,43,44,45,46.
One typical REE for alumina formers is that the RE addition tends to decrease their oxidation rates. For example, a decrease in the alumina TGO growth rate has been observed in RE-implanted8,9,15,29 and RE-alloyed8,12,15,38 β-NiAl. This aspect of REE have been interpreted by two popular models. One is a so-called poisoned interface model (PIM)47, which applies well to interpretation of the REE on growth of the cation-diffusion oxide (e.g. chromia48) scale. In the PIM, RE atoms are proposed to segregate to the scale/metal interface, pinning the climb of misfit dislocations there required for the scale growth. The other is the grain boundary (GB) segregation model. It indicates that the RE ions which incorporate into the growing alumina normally tend to segregate to the oxide GBs, where they exert the REE through either “site blocking” − blocking the fast diffusion paths for Al3+3,6,15, or a “swamping-out” mechanism49, in which the isovalent segregants (e.g. Y3+) suppress segregation of other divalent and tetravalent cations (e.g., Ca2+ and Si4+50). The aliovalent cation segregation in alumina can enhance the GB diffusivity of aluminum cations () and oxygen anions () by increasing the number of anion and cation vacancies. More recently, some investigators attributed the REE on the alumina growth to a modification of the electronic structure of alumina with GB donor and acceptor states to the extent that Al ionization at the interface is decreased30,36.
The RE segregation model has been supported by many experimental observations of RE segregation at the alumina GBs by means of TEM in a combination of X-ray energy-dispersve spectroscopy (EDS)3,4,6,7,10,12,15,17,32. In addition to the GB segregation, RE has been found to occur as RExOy particles on the alumina scale surface15,17. On its basis, a dynamic segregation theory (DST) has been proposed, in which the RE segregants at the GBs are not static; they can transport outward along the GBs driven by the oxygen potential (i.e., oxygen chemical activity) gradient across a growing alumina scale and their high affinity for oxygen15.
Many literatures15,17,20,24,25,28 also reported that addition of the RExOy dispersions in an alloy plays a similar role as RE in decreasing the oxidation rate of alumina-forming metals. This effectiveness REE of the RExOy on the alumina TGO growth is firstly attributed to the dissolution of RExOy which is proposed to occur under the oxygen pressure gradient across the metal-oxide-gas system15,24, producing RE atoms to segregate to the oxide/metal interface. However, there has been no experimental evidence for such dissolution (or dissociation) and the latter also appears to have no thermodynamic justification, because RExOy (more stable than Al2O3) have a very high thermal stability. The oxygen pressure in the metals, which decreases from the dissociation pressure of Al2O3 at the interface to some low values (depending on the oxygen solubility and diffusivity) at some distance from the interface, is normally not low enough to drive the RExOy dispersions to dissociate and release RE atoms which can segregate to the interface. In view of this, the concerned REE exerted by RExOy appears not be explained fully by DST, although it cannot be explained appropriately by PIM. A much more likelihood that the RExOy dispersions exert the REE is associated with the dissolution of RExOy dispersions upon incorporation into the alumina scales, as described simply in3,6. This raises a question on how the oxide dispersions enter the alumina TGO.
Recently, an interpretation on the effect of the ceria particle dispersion on the growth process of alumina scale on an alumina-forming aluminide was proposed32. It highlights two points. First, the incorporation of the ceria particles into the alumina TGO results from inward growth of the inner part of the alumina in α phase. Second, the ceria particles do not exhibit the REE until they have been incorporated into the alumina scale, where they may dissolve to some extent to produce the cerium ions that can segregate to the alumina GBs and thus suppress the outward diffusion of Al cations along the short-circuit paths for the TGO growth. However, this interpretation is still lack of sufficient evidences. We further characterized the microstructure of the alumina TGO on the ceria-dispersed aluminide and traced cerium either in its elementary or oxide form from the metal to the TGO. There are new observations: (i) no dissolution of original CeO2 particles in the metal during oxidation, (ii) the identification that the TGO scale is composed of the outward growing γ-Al2O3 and inward growing α-Al2O3 and the CeO2 particles in the metal can be swept over by inward growing α-Al2O3, (iii) the detection of cerium ions segregated to the GBs of the inner α-Al2O3 and (iv) precipitation of novel Ce2O3 particles along the twin boundaries in the outer γ-Al2O3. On these bases, we propose a model in the present work, which shows a scenario of a dynamic evolution of the ceria particles in the metal during oxidation. The model is helpful for better understanding of not only the REE on the alumina growth on the ceria-dispersed aluminide but also the concerned REE of the RExOy dispersions in other alumina formers. In addition, it is useful for getting insight into the alumina growth on the metals alloyed with REs, which can be preferentially oxidized into RExOy particles in the metals because the alloyed amounts of REs normally exceed their low solubility limits1,2,3,7,23,25,26,51.
Ceria dispersion in aluminide before oxidation
The CeO2 particles used, which have a typical CaF2-type crystal structure (space group:) with the lattice spacing of d(200) = 2.7 Å and d(111) = 3.1 Å on a basis of HRTEM investigation and FFT diffraction (see supplementary material 1), are in a size range of 15–30 nm. The particles were co-deposited with Ni, forming a ~35 μm-thick Ni-based composite film, in which the CeO2 particles with the content of 3.5 wt.% are in general uniformly distributed, as viewed previously by using SEM37. After aluminizing, the Ni-CeO2 composite film was converted into a ~43 μm-thick alumina-forming δ-Ni2Al328. The CeO2 particles are uniformly dispersed throughout the thickness of the aluminide on a basis of the electron probe microanalysis (EPMA)37.
Ceria dispersion in aluminide and its evolution in alumina scale after oxidation
Figure 1 shows the cross-sectioned aluminide for 30 min oxidation at 1100 °C. The aluminide forms an alumina scale. The inward growth of the alumina, as suggested by the non-planar interface, leads the metal to be either partially (as indicated by 1) or fully (as indicated by 2) enclosed by the oxide. The alumina scale viewed under TEM as seen in Fig. 2(a) displays a double-layered structure. The outer needle-like platelets, which exhibit a high density of lamellar nanotwins with coherent boundaries when tilted to the  zone axis, are γ-Al2O3 as identified by HRTEM image and the corresponding SAED pattern in Fig. 2(b). Similar lamellar-twined structure has been observed in deformed fcc γ-grains of a single-phased austenitic steels52 and in Au nanocrystal-seeded Si and Ge nanowires53. This suggests that the growth of γ-Al2O3 platelets is controlled by outward diffusion of aluminum cations along the twin boundaries (TBs) in the  orientation. The oxide of the inner layer is α-Al2O3 as unveiled in Fig. 2(c). Between the γ-Al2O3 layer and the α-Al2O3 layer appears a γ- and α-mixed area as seen in Fig. 3. The γ- and α-Al2O3 grains are the smallest in the alumina scale and Ce-rich oxide nanoparticles (see the Ce X-ray mapping) can be sometimes observed. Similar Ce-rich oxide particles occur in the inner α-Al2O3 layer. They are CeO2 as identified in Fig. 4(a), displaying the shape and CaF2-type crystal structure similar to the original CeO2. The outer highly-twined γ-Al2O3 layer is also doped cerium-rich oxide particles, which as circled in the TEM BF image in Fig. 4(b) are seemingly elongated along the twinning orientation, with respect to the particle shape of the original CeO2. The 3.3 Å lattice spacing of both and planes and (Mn0.5Fe0.5)2O3 (space group:)-similar structure (see HRTEM image and FFT diffraction in Fig. 4(b)) ascertain the particles as new Ce2O3 rather than the original CeO2.
The aluminide has been degraded from δ-Ni2Al3 into β-NiAl due to the aluminum consumption by oxidation, as shown in the inserted SAED pattern in Fig. 5(a). The aluminide contains the nano-dispersions, which have been characterized to be original CeO2. No Ce was acquired around the CeO2 particles by the EDS detector with an incident beam spot size of 1.5 nm. Figure 5(b) shows an EDS result of a specific spot between two close CeO2 particles at the GBs, showing no acquisition of Ce atoms there.
The particles of the CeO2, as inert oxide in the metal, actually act as the immobile markers for the direction of the alumina growth. They occur in the fine-grained γ- and α-mixed area (Fig. 3), suggesting that the area corresponds to the surface zone of the original aluminide. The CeO2 particles in the α-Al2O3 layer arises from the inward growth of the oxide. To further clarify this, the alumina scale formed only for 5 min has been observed. A CeO2 particle which has been swept over by inward growing alumina is clearly seen in Fig. 6. In contrast, the Ce2O3 particles in the outer γ-Al2O3 layer should be newly precipitated. They can form, suggesting that there exist sufficient Ce cations which can be migrated from the inner α-Al2O3 layer. The larger-sized ions as the Ce cations here doped in the alumina TGOs are easily segregated to and then migrate outward along the GBs15,17. As shown in Fig. 7, the Ce segregation at the α-Al2O3 GBs can be clearly seen by using HADDF-STEM. The HADDF image presents the Ce segregated GBs presents as the lines with a light contrast similar to that of the CeO2 particles (as arrowed in the BF image), because Ce has a higher atomic number than Al. The EDS analysis indicates the GBs containing a mean content of ~0.4 at.% Ce. The Ce at the oxide GBs unlikely originates from its atoms in the aluminide, because the latter have not been acquired in the metal (Fig. 5). It convincingly arises from the segregation of cerium cations, produced by partial dissolution of the CeO2 particles incorporated in the α-Al2O3 layer. The dissolved Ce cations also experience the charge transformation from tetravalent to trivalent in the alumina scale, on a basis of the precipitation of Ce2O3 rather than original CeO2.
In sum, the TEM work presents several observations: (i) the aluminide during oxidation forms an alumina scale being composed of an inner α-Al2O3 layer and an outer γ-Al2O3 layer; (ii) the CeO2 dispersions are incorporated into the α-Al2O3 layer as the result of its inward growth; (iii) Ce ions segregates to the alumina GBs and (iv) novel Ce2O3 particles are precipitated in the near surface of the γ-Al2O3 platelets.
The precipitation of new Ce2O3 in the outer γ-Al2O3 platelets demonstrates a series of evolution of the original CeO2 particles after they have been incorporated into the growing alumina, including their partial dissolution (since no evidence for such dissolution could be acquired in the aluminide (Fig. 5)), tetravalent-to-trivalent charge transformation and outward migration of the dissolved Ce cations. To unveil the dynamic evolution of the CeO2 dispersion in the aluminide, a model is schematically illustrated in Fig. 8 based on the TEM observations and interpreted below. The highly-twined γ-Al2O3 grains grow outward quickly on the aluminide at the onset of oxidation and α-Al2O3 nucleates at the γ/aluminide interface soon after if not simultaneously (Step I). The γ-Al2O3 grains grow into needle-like platelets in the direction preferentially aligned with the  twinning orientation; in the meantime, the initially-formed α-Al2O3 grains gradually spread laterally and inward, sweeping over the CeO2 particles in the aluminide (Step II). α-Al2O3 exhibits an n-type behavior with the principal defect of oxygen vacancy or free electron e′ (Vink-Kröger’s notation). The CeO2 partially dissolves into the α-Al2O3 lattice through the reaction
where represents the quaternary-charged Ce cations and lattice oxygen. The in the n-type oxide lattice can then trap an electron and transform to trivalent-charged cation through the reaction
The reduction of CeO2 to Ce2O3 has been reported in the high temperature sintering of fine CeO2 particles54,55. The (1.02 Å with the coordination number of 656) is larger than (0.87 Å with the same coordination number) in the ion size. Larger in the α-Al2O3 grains yields higher lattice misfit microstrain, which drives the trivalent cations to segregate to the α-Al2O3 GBs (Step III). Then, the segregated cations migrate from the α-Al2O3 layer to the γ-Al2O3 layer along the GBs in the α-Al2O3 and the lamellar TBs in the γ-Al2O3 platelets, under the driving force of the oxygen potential gradient across the oxide15,17. γ-Al2O3 is a p-type oxide with the principal defect of Al vacancy and electron hole . The TBs, although they are coherent, contain steps and kinks which can serve as the sinks for vacancy (as having been reported in57,58) like here. Steps and kinks in the lamellar TBs in the γ-Al2O3 can trap . Once and are both oversaturated there, Ce2O3 is precipitated (Step IV) through the reaction below,
The cations prefer to diffuse outward along the TBs, causing the Ce2O3 to be precipitated and elongated in the growth direction of the γ-Al2O3 platelets (Fig. 4(b)). The precipitates are easily observed in the γ-Al2O3 layer near the surface, because of higher there which promotes the precipitation reaction.
The oxidation kinetics of alumina formers during 750–1200 °C is highly correlated with the diffusion of Al cations and O anions along the alumina GBs30. A decrease of by the RE segregations to alumina GBs has been proposed to be the reason why the RE- and RExOy−doped alumina formers have a lower oxidation rate3,4,6,7,10,12,15,17,32. As illustrated in Fig. 8, in the alumina layer here should be decreased when the Ce segregants outward migrate along the GBs in α-Al2O3 layer and the TBs in the γ-Al2O3 layer.
The model in Fig. 8 strongly suggests that the Ce segregants occur only when the CeO2 dispersoids in the aluminide have been swept over by the inward moving alumina/metals interface. In other word, the REE on the alumina growth for the CeO2 dispersion in the aluminide is intrinsically pertinent to the incorporation of the oxide particles into the alumina scale by its inward growth. This may be generalized to the REE of the other RExOy dispersions on alumina TGO growth. In addition, many alumina formers are alloyed with a RE instead of its oxide. Theoretically, it is possible that the RE in an alumina-forming metal at and below the equilibrium partial oxygen pressure of Al2O3/metal at the interface can be internally oxidized to form RExOy. For example, Hf in an alumina-forming CoCrAlHf was internally oxidized into fine spherical HfO2 particles in an Al2O3/CoAl pack at 1200 °C2. The particle sizes of the formed RExOy highly depend on the synergistic effect of several factors, e.g., RE amounts and solubilities in metals, metal compositions and microstructures, alloying and oxidation temperatures25,38,51. The particles have not been highlighted previously, plausibly because they are sometimes as small as nano-sized particles. Because of a very low solubility limit in metals (e.g., only 0.01 wt % for Y in the FeCrAl alloys51), a RE overalloying is hard to avoid. Thus, RE-rich precipitates occur in metals1,2,3,7,25,26,51. They can be in-situ internally oxidized to form RExOy particles at the front of oxidation. The RExOy particles in the RE-alloyed metals, no matter whether they are formed by diffusional and non-diffusional oxidation of RE solute atoms and RE-rich precipitates, respectively, would not have the relative REE until they have been swept by the inward growing alumina.
In summary, the REE for the CeO2 particle dispersion in the nickel aluminide is firstly correlated with the inward growth of the inner α-Al2O3 in the TGO scale. The CeO2 dispersion, after being swept by the inward-growing α-Al2O3, partially dissolves producing tetravalent . They then transfer to trivalent by capturing electrons and segregate to the GBs in the n-type α-Al2O3. The cations migrate outward along the GBs there and the TBs in the outer γ-Al2O3, finally precipitating Ce2O3 near the surface. The segregation and migration of the Ce cations along the planar defects would obstruct the diffusion of the Al cations for the growth of the alumina TGO. The explanation may be generalized to the related REE of other RExOy dispersions in the alumina-forming metals. For the alumina formers alloyed with a RE, the RE, in the form of either solute atoms or RE-rich precipitates, is plausibly internally oxidized into RExOy particles below the TGO/metal interface. They may then affect the alumina growth in a manner to the CeO2 dispersion in the aluminide.
The CeO2 particles with a purity of 99.5%, a commercial product by Alfa Aesar company, were introduced to a nickel aluminide by using a two-step method28,37. First, the pure Ni samples with dimensions of 15×10×2 mm, after being abraded to a final 800 grit SiC paper, were electrodeposited with a Ni-CeO2 composite film from the CeO2-loaded nickel sulfate bath (150 g/l NiSO4·6H2O, 120 g/l C6H5Na3O7·2H2O, 12 g/l NaCl, 35 g/l H3BO3). A mechanical agitation was maintained to mitigate the particle agglomeration and sedimentation during electrodeposition, as illustrated by the setup59. Second, the samples were aluminized at 620 °C for 5 h using a halide-activated pack-cementation in a powder mixture of Al (particles size: ~75 μm) + 55 wt.% Al2O3 (~75 μm) + 5 wt.% NH4Cl in an Ar (purity: 99.99%) atmosphere. The characteristics of inward growth of the aluminide at the cementation temperature caused the CeO2 in the electrodeposited film to be trapped, forming a ceria-dispersed nickel aluminide coating on the sample surface. After being ultrasonically cleaned in acetone, the aluminized samples were ready for oxidation.
The samples were not placed into a muffle furnace for oxidation until it was heated up to 1100 °C. The ceria–dispersed nickel aluminide after oxidation were cross-sectioned for the scanning electron microscopy (SEM) investigations and then ion sliced into thin foils by using techniques detailed elsewhere32, with the transparent areas desired for the TEM investigations under a JEOL 2100F TEM at 200 kV accelerating voltage, by using the techniques of bright field (BF) imaging and high resolution TEM (HRTEM) imaging and fast Fourier transformation (FFT) diffraction, scanning TEM (STEM) imaging as well as high-angle annular detector dark-field STEM (HADDF-STEM) imaging.
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The work is supported by National Natural Science Foundation of China (NSFC, project Grant No. 51471178). The author, XP, thanks Dr. P. Y. Hou at Lawrence Berkeley Natl. Lab. USA, for her helpful comments.
The authors declare no competing financial interests.
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Wang, X., Peng, X., Tan, X. et al. The reactive element effect of ceria particle dispersion on alumina growth: A model based on microstructural observations. Sci Rep 6, 29593 (2016). https://doi.org/10.1038/srep29593
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