Development of Attribute Maps for Grain-Boundary Transport in Doped Alumina

The growth of protective alumina scales in Al203-fonning alloys can be affected by the addition of reactive elements, such as Hf*"^, which has been considered one of the most effective dopants to slow down the scale growth rate. While a number of theories concerning the "reactive element effect" have been proposed, a full explanation of this phenomenon is not yet available. The overall objective for this study was to conduct a systematic series of model experiments in order to elucidate the effect of Hf02 on oxygen grain boundary transport in alxmiina. The key questions to focus in the current work are: How do the doping levels of HfOi / oxidizing temperature affect oxygen grain boundary diffusion in alumina? First part of this work investigated the effect of doping levels of HfDa on oxygen grain boundary transport in alumina, which contains uniformly distributed Ni metallic particles. The doping levels spanned the solubility limit ranging from lOOppm to 2000ppm. The plot of the ratio kundoped/kdoped (grain-size corrected) as a function of dopant level clearly shows two behavior regimes: namely a regime I that encompasses doping levels below and near the solubility limit and a regime U where second-phase Hf02 particles were well present in the microstructure. A clearer understanding of the influence of Hf02 doping on the transport behavior can be achieved by plotting the data with respect to the fractional grain boundary coverage (/), as opposed to overall Hf02 content. The linear relationship can be rationalized with a site-blocking model, in which the Hf*"^ ions obstruct the diffusive paths at the grain boundary. 1 The second part of the work is focused on the temperature dependence of the oxidation kinetics in Hf02 doped AI2O3. The activation energy and rate constant ratio plot from oiu: work and alloys studies indicated that multiple diffusion mechanisms might be operative at graia boundaries owing to boimdary transitions that modify local structure and chemistry. Resutls of ARM characterization of samples oxidized at 1150°C have revealed that boundary structures that differ from those observed in samples oxidized at 1250°C and 1400°C. Specifically, new types of boundary structures are present at the lower temperature that are more atomic rough structure, and exhibit high energy facet planes such as {0 0 0 6} and {2-1-13} different from relative low energy facet plane such as {2 -1 -1 0} and {1 0-1 2} for samples oxidized at higher temperature. These high-energy facet planes maybe occupied with high Hf segregation level and resulted in a much better oxidation resistance. The current results highhght the significant role of complexion fransition in the oxidation area, which is fraditionally being neglected. Chapter 1. Background 1.1. The Growth of Alumina Scales Gas turbines, which are widely used in aircraft, marine and power transmission area, have progressed greatly over the last several decades based on extensive research efforts to increase the operating temperature of the engine and therefore, increase the efficiency of the engine with the development of internally cooled components and thermal barrier coatings. [2, 3, 84, 86] However, as the application temperature rises, ^ oxidation resistance will be a primary degradation mode, especially for turbine blades, which experiences the most severe environment such rapid temperature transients, oxidizing gases, contaminants such as chlorides and sulphates. Among various commercial alloys. Nickel-based superalloys are the widely used turbine blade materials with an exceptional combination of high temperature strength, toughness, and resistance to degradation in corrosive or oxidizing environments, which can tolerate temperature as high as 1200°C that is approximately 90% of the melting point of the material. [84, 87] Besides the FCC Nickel as the major superalloy constituent, there is a combination of five to ten other elements alloyed with nickel to achieve a significant improvement in properties such as creep, oxidation and so on, as listed in Figure 1.1. 3 In order to realize greater efficiency, the next generation of gas turbine systems will operated at increasingly higher temperature, which is even above the meltiag point of superalloys from which they are comprised. With the deposition of a multiplayer and multicbmponent thermal barrier coatiugs (TBS), the turbine blades can survive up to 40,000 hours before failing. [3, 84, 85, 88] Figure 1.2 shows a typical cross section view of turbine blades with TBC coating. The topmost layer is typically a 6~8wt% yttria stabilized Zirconia (YSZ), which act mainly as a heat shield to provide major reduction in the surface temperature. (100°C ~ 300°C) Moreover, YSZ has a lower density and well-matched thermal expansion coefficient with the xmderlying alloy, which resulted in a lighter weight and minimize the stress build-up during the thermal expansion mismatch. However, YSZ is transparent to oxygen diffusion due to its high oxygen vacancies and oxidation from xmderlying superalloys resulted in fast growing Ni-rich oxides, which are not thermodynamically stable with YSZ [86] and hence cause failure of the coating. Therefore, a bond coat alloy (MCrAlY, M=Ni, Co) is usually placed between YSZ and superalloys. During the service, oxidation of the bond coat results in the formation of thin alumina layer between the bond coat and YSZ, which can provide oxidation resistance by slowing fiirther oxygen diffusion through this thermal grown alumina. And compared with other protective oxides such as Cr203 or SiOi, alumina is widely used due to its relative ease to form a continuous layer, its slow growth rate and long-term chemical stability since Cr203 has volatility issues at temperature higher than 800°C and Si02 has volatility issues with H2O. [2] However, with the alumina growing thicker and thicker, various failure mechanisms could initiate such as stress-induced delamination of the thick alumina due to thermal expansion mismatches and reduced interface adhesion between alumina and bond coat due to the presence of interfacial voids, which will expose the substrate alloys directly to oxidizing environment. [2, 3, 85, 89] Since there is no practical strategy for preventing the high-temperature oxidation, it is essential to reduce the alumina growth rate to maximize the use temperature and extend imderlying superalloy's lifetime. Considering the growth of the alumina protective layer is a diffusion controUed-process, understanding the transport mechanisms of Al and O in the oxide layer will be essential. Various methods have been used to elucidate the growth mechanism of this thermal grown alumina. One way to investigate the mechanism is by placing inner markers on the non-oxidized metal surface. After oxidation, the location of the inner markers could indicate the alumina is formed by either inward O diffusion or outward Al diffusion, or both. Ramanarayan et al. [90] used tiny droplets of Pt markers in FeCrAl alloys. The results showed that the Pt marker was placed at the gas-alumina interface, which strongly indicated that the growth mechanism is occixrred by O inward diffusion. While results from Mrowec et al. [91] by using nano-size Au markers showed that the alumina was grown by simultaneously O inward and Al outward diffusion. Moreover, the rehability of these marker experiments has been questioned to have misleading results. In Yoxmg et al [92] study, Pt Markers remained on the surface of the NiAl alloys after oxidation at 900°C suggesting inward oxygen transport, while O tracer study in the same alloy showed the contribution of Al outward diffusion cannot be neglected. Another way to study the growth mechaism of alumina scales is by using Secondary Ion Mass Spectrometry (SIMS) techniques combined with ^^O tracer experiments. The basic principle can be described below: The alloys are first oxidized in the '^O atmosphere to form a thin layer of ^^O-containing alumina scales. Then the alloy was oxidized again in an ^^0-rich gas. By using SIMS, the depth profile of ^^O concentrations as a function of locations through the alumina and alloy systems will give information on the alumina growth mechanism. Various transport mechaims have been proposed by Reddy et al. [93] and Quadakkers et al. [94] based on the distribution of the isotopes. It is reasonable to speculate that four possible transport mechanisms fi-om different ^^O concentration profiles. If outward Al diffusion is the dominant mechanism, the ^^0-containing alumina will be formed at the gas-oxide interface and the tracer concentration will be a step function. If inward O diffusion is ID the primary mechanism through pores or microcracks, then the 0-containing oxide will be formed at the oxide-alloy interface and the tracer concentration will also be a step function. Moreover, if the alumina is mainly grown through grain boundary diffusion, most of the ^^0-containing alumina will be formed at the oxide-alloy interface while minor amount will be formed at the gas-^^0-containing oxide interface due to exchange of ^^O with ^^O at the grain boundaries. If the diffusion mechanism is a combination of outward Al and inward O grain boundary transport, the ^^0-containing oxide will be formed at both gas-oxide and oxide-alloy interface. It should be pointed out that the dominant growth mechanisms could be changed by varying oxidation temperatures. Mitchell et al [95] found the Al outward diffusion is the primary mechanism for NiAl oxidized at 1100°C, while the contribution of the O inward diffusion becomes more significant with increasing oxidation time and temperature. Similar trends have been observed by Young et al. [96] that NiAl oxidized at 900°C indicated a mainly Al outward diffusion, while increased inward O diffusion at 1150°C. The third method is based on the microstructural evidence dxiring the growth of alumina scales proposed by Tolpygo and Clarke. [18, 97] The outward Al diffusion can be analyzed by measuring the oxide ridges growing on the alumina surface, while the inward O diffusion was inferred from the columnar structure of the alumina grains. The schematic illustration of the experiments is shown in the Figure 1.3. The basic procedures can be described below: The alloy was first oxidized to form a thin and adherent alumiaa layer. Then the top surface was mechanically polished to remove the oxide layer at a small angle to the sample surface. After oxidation again, the surface morphology can be analyzed by using SEM. The newly formed oxide ridges at the alumina grain boundaries on the top surface were grown mainly by Al outward grain boundary diffusion. By knowing the outward diffusion flux, the thickening of the oxide by inward O diffusion can be obtained from weight-gain measurements. Therefore, the outward Al flux can be quantified as a function of oxide thickness and information about the ratio of outward diffusion of Al to inward diffusion of O can be obtained. IIA IIIA IVB ^x-Element B 0.097 0.077. -Atomic Radius (nm) ■ rvA VA VIA VIIA VIIIAVIIIAVIIIA ^^■0132 Cr 0.125 Fe 0.124 Co 0.125 Ni 0.125 Y 0.181 0.15l'"ffllH Mo 0.136 Ru 0.134 Hf 0.159 , Ta .0.147 w 0.137 Re 0.138 Y' former | | Minor alloying additions | | y former Figure 1.1: Normal alloying elements presented in Ni-based superalloys [84]. Superafloy Bond Coat ~60|.ifn Ceramic Top Coat Cooling Air Ftfm -lOO^OOum /