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Compositional effects on the high-temperature oxidation lifetime of MCrAlY type coating alloys
Corrosion Science 95 (2015) 143–151
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Compositional effects on the high-temperature oxidation lifetime of MCrAlY type coating alloys S. Salam a, P.Y. Hou b, Y.-D. Zhang a, H.-F. Wang a, C. Zhang a, Z.-G. Yang a,⇑ a b
Key Laboratory of Advanced Materials, Ministry of Education, School of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA
a r t i c l e
i n f o
Article history: Received 17 January 2015 Accepted 9 March 2015 Available online 14 March 2015 Keywords: A. Alloy B. Modeling studies C. Oxidation
a b s t r a c t The lifetimes due to chemical failure during high-temperature oxidation of Co and Ni-rich MCrAlY type coating alloys were investigated. They were found to depend critically on the oxidation rate and Al diffusivity in alloy. Re and Al additions increased the life by providing higher b-phase contents and by lowering the Al diffusivity and oxidation rates. The effects of Re were more pronounced in the Co-rich than the Ni-rich alloys. The observed difference was a result of different alloy phases, with a microstructure of c, b, and r in the Co-rich alloys and c, b, and a in the Ni-rich alloys. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction Increase in the operating temperatures of modern gas turbines for improved efficiency have resulted in higher thermal–mechanical fatigue and oxidation of the turbines’ superalloy components [1,2]. To combat these problems, thermal barrier coatings (TBCs), consisting of a ceramic top coat (TC) and a metallic bond coat (BC), are usually applied to the substrates [2,3]. MCrAlY alloys, where M = Co, Ni or their combination, are one of the commonly used BC material. A thin surface layer of a slow-growing Al2O3 is developed between the TC and the BC at high temperatures by the preferential oxidation of Al in the MCrAlY [4–6]. This leads to the consumption of Al such that a subscale region is formed ahead of the Al2O3 layer, within which Al is depleted [7–9]. The consequent compositional changes in the alloy can give rise to microstructural changes where prolonged exposures may cause TBC failure [10,11]. TBCs applied in modern gas turbines, operating at a bond coat temperature of 900–1000 °C, require a lifetime of 25,000 h or more [12]. The TBC life is limited by a number of potential mechanisms including thermal fatigue failure, thermal aging failure, and Al depletion failure [12,13]. Out of these three, the Al depletion failure is considered as chemical breakdown, which occurs when the Al content in the alloy reaches a critically low value that can no longer sustain the continued growth of the protective scale. Consequently, less protective oxides, like fast-growing Ni(Co, Cr)spinels, are formed [14]. Experimental determination of the TBC ⇑ Corresponding author. Tel.: +86 1062783848; fax: +86 1062771160. E-mail address: [email protected] (Z.-G. Yang). http://dx.doi.org/10.1016/j.corsci.2015.03.011 0010-938X/Ó 2015 Elsevier Ltd. All rights reserved.
lifetime is not only tedious but also expensive. Instead, numerical analysis is preferred, which describes the alloy phase equilibria from calculated concentration changes in the alloy during hightemperature oxidation [15–18]. Often the time taken for an Al2O3 forming alloy to be completely depleted of its b phase is accepted as its chemical lifetime [19,20]. Recently, computer aided technologies for the modeling of phase diagrams and diffusion processes, like Thermo-Calc and DICTRA, have shown promising results to calculate the concentration changes underneath the scale during oxidation [21–23]. Microstructural changes arising from compositional variations could affect key bond coat properties, e.g. oxidation rate and diffusivity, which would affect the service life of the alloy. Oxidation rate of the bond coat alloy is clearly an important parameter, since the alloy with a higher rate would result in a more rapid consumption of the scale forming element. Diffusion in the alloy may also be important since it can deplete the alloy of essential constituents, such as Al, due to coating/substrate interdiffusion [11,24]. Existing experimental lifetime values are usually reported for thermal cycle failures [25,26], which are much lower than chemical lifetimes. Published chemical lifetime data vary substantially among different TBC systems [19,20]. These differences are attributed to parameters such as bond coat compositions, coating thicknesses, substrate compositions, operating temperatures, and deposition methods. Many researchers have attempted to improve the alloy’s oxidation resistance and lower the consumption of the scale forming element by improving the composition of the bond coat alloy [27–32]. However, the effects of microstructural phases independent of composition have received relatively little attention.
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In this study, CoNiCrAlReY alloys with different compositions were chosen. Prior works on these alloys had examined the effects of Al concentration and Re addition on their oxidation behavior [27], and had modeled the alloy phases resulted from these compositional changes [33]. The aim here is to investigate the influence of such compositional changes on the alloys’ chemical lifetime. Furthermore, in an effort to differentiate the effects of alloy phases from its composition, a number of Co- or Ni-rich MCrAlReY alloys were prepared, which provided alloys with different phase fractions of c + b with or without a or r precipitates. The chemical lifetimes of these alloys were calculated using Thermocalc and DICTRA. The lifetimes were then correlated with the effects of phases and compositions on the alloys’ oxidation and Al diffusion rates.
Fig. 1. Use of EPMA line scans to collect average concentration profiles over different positions along the diffusion direction in the alloy.
diffraction (XRD) (Rigaku/SmartLab) using Cu ka radiations at 40 kV.
2. Experimental method 2.1. Experimental work
2.2. Modeling procedure
Co and Ni-rich alloys were melted in a vacuum electric arc furnace (SK-II non-consumable vacuum arc furnace) from weighted percentage of pure elements in the form of high-purity (99.99%) metallic ingots obtained from Beijing Jiaming Platinum Industry of Non-Ferrous Metals Ltd. Co and Ni-rich alloys were melted to form button-shaped ingots of 100 diameter weighing 80 g each. Following melting, the ingots were homogenized in vacuum at 1150 °C for 24 h followed by the air cooling to room temperature. Actual compositions of the experimental alloys, determined with inductively coupled plasma mass spectrometer (ICP-MS) (Thermo Scientific™ iCAP™ Qc ICP-MS), are listed in Table 1. These alloys are given short names for simplicity, e.g. CoRe-8 means Co-rich alloy containing Re with 8% Al. Small disc-shaped samples (£10 mm 1 mm) were cut using an electrical discharge machine (EDM). Samples were mechanically ground and polished down to 0.5 lm grit using diamond paste. Finished samples were placed in alumina boats, containing 5 samples each, and inserted into a box furnace (Nabertherm LHT04/17) for oxidation in the air. Alloys were oxidized at 1000 °C for 100, 250, 500 and 1000 h. Additional samples were oxidized at 950, 1050, 1100, 1150, 1200 and 1250 °C for 100 h. For cross-sectional examination, oxidized samples were mounted in epoxy in a vertical orientation and then polished. The concentration profiles in subscale region were measured with an electron probe micro-analyser (EPMA) (Jeol JXA-8230) coupled with wavelength dispersive spectrometer (WDS). In order to accurately obtain the average concentration of different elements along their diffusion trajectory, a number of line scans were preformed parallel to the oxide/alloy interface as shown in Fig. 1. The length of each line was 100 lm, and the lines were spaced 2 lm apart. As seen from Fig. 1, these line scans pass through different phases; the elemental concentration at the location of each line was averaged over all the phases. Microstructures and phase compositions of the alloys were determined with a scanning electron microscope (SEM) (Jeol JSM-7001F). Alloy phases were identified with X-ray
Al consumption and its distribution in the alloys was modeled with DICTRA and Thermo-Calc using the thermodynamic and kinetic parameters obtained from the databases TTNI8 [34] and mob2 respectively. A continuous matrix (c phase) was considered containing b, and other precipitate phases dispersed within. The model assumes that diffusion takes place in the c matrix. The precipitate phases affect the diffusion process by their chemical composition and volume-fraction. Al was removed from the alloy with a flux proportional to the oxidation rate (kp) [35,36] using a hypothetical diffusion couple between an imaginary vacuum and the alloy. The vacuum acted as a sink for the Al leaving the alloy during oxidation. For a small time step Dt, Al in the alloy is removed through the diffusion couple, and the diffusion equations for the matrix phase are solved. New compositions are obtained at each point in the alloy assuming local equilibrium. For the next time step Dt, DICTRA solves the diffusion equations gain for the newer composition, etc. Using this simple modeling approach, the Al concentration and its distribution in the alloys was calculated as a function of oxidation time. Corresponding phase fractions were obtained from concentration changes within the alloy using calculated phase diagrams. Coating thickness is an important parameter while determining the total amount of b in the alloy. A typical bond coat thickness of 200 lm [19,20] was used to carry out the calculation. The total amount of the b phase in the alloy was continually calculated; the time taken for a complete b dissolution provides the chemical lifetime of the alloy.
Table 1 Composition of experimental alloys, determined by ICP-MS. Alloy ID
CoRe-8 CoRe-10 CoRe-12 Co-10 NiRe-10 Ni-10
Composition in wt.% Co
Ni
Cr
Al
Re
Y
Bal. Bal. Bal. Bal. 8.1 ± 0.4 8.2 ± 0.3
32.2 ± 0.2 31.9 ± 0.5 32.1 ± 0.3 31.9 ± 0.2 Bal. Bal.
21.3 ± 0.9 21.1 ± 0.7 20.9 ± 0.5 21.3 ± 0.3 20.9 ± 0.8 21.2 ± 0.1
8.1 ± 0.4 9.9 ± 0.6 11.8 ± 0.7 10.1 ± 0.3 9.9 ± 0.2 10.0 ± 0.2
3.4 ± 0.1 3.5 ± 0.3 3.4 ± 0.4 0 3.5 ± 0.1 0
0.13 0.15 0.12 0.11 0.12 0.13
3. Results 3.1. Microstructural characteristics Fig. 2 shows typical microstructures of the experimental alloys after they were heat-treated at 1000 °C for 500 h; corresponding XRD results are shown in Fig. 3. All alloys consisted mainly of the c and b phases, which appeared as light and dark areas respectively. c phase appeared as a continuous matrix phase. The microstructure of alloys without any Re addition (Co-10 and Ni-10) consisted of c and b phases only. Many white particles were observed at the c and b phase boundaries of the Re containing alloys. XRD analysis identified these particles as r phase in the Co-rich alloys and a phase in the Ni-rich alloys. No Ni–Y intermetallic phases were detected using XRD, because of the relatively low Y content in these alloys. Composition of several white particles (a or r) were analyzed with WDS, and their average concentrations are given in Table 2.
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Fig. 2. Backscattered electron images (BEI) of experimental alloys after heat-treatment at 1000 °C for 500 h; (a) CoRe-8, (b) CoRe-10, (c) CoRe-12, (d) Co-10, (e) NiRe-10, (f) Ni-10.
Fig. 3. XRD patterns of heat-treated Co-rich and Ni-rich alloys with and without Re. Alloys were heat-treated at 1000 °C for 500 h.
Both types of particles were enriched with Cr and Re, whose concentrations were higher in a than in r. The r particles also contained a significant amount of Co; no Y was detected in r and the Al concentration in them was very low. The a phase, on the other hand, contained more Al and a notable amount of Y.
The influence of heat-treatment temperature on alloy phases is demonstrated in Fig. 4. The amount of b and a/r phases decreased with increasing temperature, and the a/r phases eventually dissolved, due to increased solubility of Cr and Re in c at higher temperatures [37]. A binary c + b microstructure was observed
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Table 2 WDS analysis showing the average composition of white particles in Fig. 2. White particles
r phase a phase
Composition in wt.% Co
Ni
Cr
Al
Re
Y
26.8 ± 0.9 2.5 ± 1.2
7.2 ± 0.4 5.9 ± 2.6
41.4 ± 0.3 56.6 ± 1.0
0.3 ± 0.1 1.7 ± 0.2
24.3 ± 0.6 30.4 ± 1.2
0 2.96 ± 0.4
for the CoRe-10 and NiRe-10 alloys heat-treated at P 1150 °C and 1200 °C respectively. The relative amount of different phases in each alloy was estimated using image analysis. A detailed method of obtaining the weighted percentage of phases from their measured surface fraction is given in an earlier paper [33]. The calculated phase amounts are shown as a function of heat-treatment temperature in Fig. 5. Co and Ni-rich alloys had similar phase content at all tested temperatures that consisted of c and b only. All six alloys showed decreasing fractions of b and r or a with increasing temperature. For a given temperature, higher amounts of b were present in alloys with more Al (compare CoRe-8, CoRe-10, and CoRe-12). The r phase precipitated in alloys with Al P 10% and increased with increasing Al content in the alloy. Re addition not only resulted in the precipitation of r or a, it also increased the alloy’s b content. This effect was slightly more pronounced in the Ni-rich than in Corich alloy. 3.2. Oxidation and diffusion All investigated alloys underwent uniform oxidation. XRD analysis through each alloy’s surface confirmed that the oxide was mainly a-Al2O3 with small amounts of Co,Ni spinels formed during the initial stages of oxidation. The oxidation rate of the alloys during the transient stage was significantly higher than in the steady-state (>10 h). A few examples of sample mass gains as a function of oxidation time are shown in Fig. 6a. A quadratic curve with substantial mass gain during the first few hours can be seen. Similar two stage oxidation behavior has been extensively reported in the literature for Al2O3 forming alloys [38,39] where initially fast
growing (Co, Ni)-spinels and transition Al2O3 form on alloy surface. A continuous a-Al2O3 does not form until at least few hours have elapsed [40,41]. Consequently, in the transient stage the slope of mass gain curves is much steeper than those seen in the steadystate. To approximate the rate constants for a quantitative comparison, the relatively parabolic section of the curve was fitted to the parabolic rate equation given by [40]:
Dm 2 ¼ kp t þ C A
ð1Þ
where Dm is the mass gain, A is the specimen’s surface area, kp is the rate constant and t is the oxidation time. Fig. 6b plots the square of the mass gains vs. the oxidation time in the second stage. Parabolic time dependency of the a-Al2O3 growth is seen for all studied alloys. The first data point of each curve in Fig. 6b represents the end of the transient stage oxidation. Approximate parabolic rate constants (kp) for the steady-state were obtained from the slopes of the lines in the parabolic plot. These kp values are indicated in the figure caption, and they are in good agreement with those reported for a-Al2O3 growth in similar alloys [42]. Fig. 6b shows that the alloys with more Al oxidized slower. The Co-rich alloys oxidized at a slightly higher rate than the Ni-rich alloys. Re addition, decreased the steady-state rate in both types of alloys, but the effect was stronger in the Co- than in Ni-rich alloys. After oxidation, polished cross-sections were examined using SEM. A dense and uniform Al2O3 scale was seen on the alloy’s surface in all the test samples. Beneath the scale, a depth containing only the c phase, identified as a b-depletion layer, was observed in all alloys. An example is shown in Fig. 7 with the CoRe-10 alloy. It is seen that the depth of the b-depletion layer
Fig. 4. Microstructure of CoRe-10 and NiRe-10 alloys after heat-treatments at different temperatures; (a–c) CoRe-10 alloy at 1100 °C, 1150 °C and 1200 °C, respectively; (d–f) NiRe-10 alloy at 1100 °C, 1150 °C and 1200 °C, respectively.
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Fig. 5. Phase content of experimental alloys, in wt.%, as a function of heat treatment temperature for (a) b, (b) c, (c) r and (d) a. Solid lines represent the Co-rich alloys and dotted lines represent the Ni-rich alloys.
increased with increasing temperature and oxidation time. Slight accumulation of the Cr–Re particles were observed at the scale/alloy interface of the alloy oxidized at 1000 °C; this behavior has been documented and explained in an earlier paper [35]. The thickness of the b-depletion layer after 100 h oxidation in CoRe-10 and NiRe-10 alloys was measured in the temperature range between 1000 and 1250 °C. These results are presented in Fig. 8, where the layer thicknesses are shown as a function of 1/T. Vertically dashed lines are used to indicate the boundary between two microstructure types, i.e. c + b at higher temperatures and c + b + r or c + b + a at lower temperatures. As the depletion thickness increases with increasing temperature, linear behavior was expected in the Arrhenius plot. This was indeed the case for the NiRe-10 alloy, but for the CoRe-10 alloy, two distinct slopes were present, suggesting that the depletion rates in the two regions were different. Within the temperature range of 1000 and 1100 °C, the microstructure consisted of c + b + r phases and the Co-rich alloy was depleted at a lower rate than in the temperature range of 1100 and 1250 °C where no r was present. The depletion rate in the Ni-rich alloy remained the same throughout the tested temperature range. The Al diffusivity in the c matrix was estimated using the analytical solution proposed by Whittle through the equation [14]:
rffiffiffiffiffiffiffiffi 1 pkc X nL ðn þ 1ÞL Ni ¼ No erfc pffiffiffiffiffiffi þ erfc pffiffiffiffiffiffi 2D n¼0 2 Dt 2 Dt
ð2Þ
where Ni is the experimentally determined interface Al concentration, No the Al concentration in the alloy, kc the oxidation rate constant in (cm2 s1), L the sample thickness, t the oxidation time and D
is the diffusion coefficient. To keep the units consistent, the measured kp was converted to kc using the equation:
kc ¼ 2
Mx VZ x
2
kp
ð3Þ
where Mx is the atomic mass of oxygen, V the volume of the oxide scale and Zx the valency of oxygen. The calculated values of DAl in the gamma matrix for each of the studied alloys are presented in Fig. 9. Data shows that Al diffused faster in the Co-rich than the Ni-rich alloy. Increasing the alloy Al content from CoRe-8 to CoRe12 decreased the Al diffusivity. Re addition also decreased the Al diffusivity; the effect was considerably stronger in the Co-rich alloy than in Ni-rich alloy.
3.3. Lifetime prediction Al concentration profiles in the alloy after different oxidation times were calculated using the method explained in the modeling section. To verify the accuracy of the modeling method, profiles for each alloy were calculated for 250 and 500 h oxidations at 1000 °C and compared with experimental values determined using EPMA; good matches were observed between the two. An example is shown in Fig. 10, where Al concentrations (solid lines from modeling and data points from EPMA measurement) in CoRe-10 and NiRe-10 alloys are plotted against the distance from the scale/alloy interface into the bulk alloy. The width of the b-depletion layer is marked as a vertical dash line, whose location was determined at the Al concentration above which b becomes stable (based on the phase diagrams calculated using Thermo-Calc). The thicknesses
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Fig. 6. Oxidation kinetics of the experimental alloys at 1000 °C; (a) mass gain as a function of oxidation time; (b) parabolic plot of (a). The slope of each curves in (b) represents the steady-state kp of the alloy. Calculated kp values in g2 cm4 s1 are shown in the legend.
of these calculated b-depletion layers were in good agreement with those observed from the sample cross-sections. The b phase in the alloy is consumed with increasing oxidation time until its content reaches zero. Lifetimes predicted by the complete consumption of each alloy’s b phase are presented in Table 3.
These modeled results show that alloys with larger b reservoirs (more Al) have longer lives; Co- and Ni-rich alloys have similar lifetimes that increase with Re addition; Re in the Co- and Ni-rich alloys increases their lifetimes by 2600 h and 1700 h respectively. Table 3 also includes the effect of coating-substrate interdiffusion on lifetime. When Al in the bond coat diffuses into the substrate, less Al is available to continue the growth of the protective Al2O3 scale. Diffusion couples between the tested alloys and a third-generation superalloy substrate (CMSX-10 with about 6% Al) were simulated with DICTRA to determine the extent of Al interdiffusion between these two. An example of the resulting diffusion profile, between CoRe-10 and CMSX-10 after 500 h at 1000 °C, is shown in Fig. 11. The vertical dash line in the figure represents the original interface between the coating alloy and the substrate. As seen from Fig. 11, the Al back-diffuses into the substrate and the Al concentration in the coating decreases. A subsequent Al depleted single c phase zone (the marked depletion layer-II on Fig. 11) is formed, and its depth increases with time. This profile indicates that Al in the single crystal substrate with no grain boundaries diffuses at a slower rate [12], while the polycrystalline coating alloys with grain boundary diffusion allowed faster diffusion in the coating. Calculated width of interdiffusion zone in the coating did not extend more than 50 lm. For the lifetime modeling, the width of this depletion layer-II at any given time was calculated and then subtracted from the coating’s original thickness. This correction undoubtedly reduces the total amount of the b content in the alloy and consequently reduces the lifetime. The corrected lifetimes and their % decrease from the oxidation-only condition without any diffusion into a substrate are shown in Table 3. The values in the table show that coating lives decrease by 12–21% due to Al back-diffusion into the substrate. The extent of this back diffusion decreases with increasing Al content in the alloy and with Re addition.
4. Discussion Minor changes to the alloy’s composition affect the bond coat properties that dictate the chemical life of the alloy. Results from this study are summarized in Table 4, which show the effects of
Fig. 7. SEM cross-sectional images of CoRe-10 alloy oxidized at (a) 1000 °C for 100 h, (b) 1000 °C for 500 h, (c) 1100 °C for 100 h and (d) 1200 °C for 100 h. A b-depletion layer can be seen beneath the Al2O3 scale under all conditions.
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Fig. 8. Temperature dependence of the b-depletion thickness in Re-containing alloys oxidized for 100 h; (a) CoRe-10, (b) NiRe-10.
Fig. 9. Al diffusivity in the experimental alloys calculated at 1000 °C.
alloy chemistry and its resulting microstructure on the oxidation rate, the Al alloy diffusion rate and the chemical lifetime of tested Co-rich and Ni-rich coating alloys. The lifetime of an alloy can be increased by adding Al or by modifying its composition to reduce the Al consumption during subsequent oxidation and interdiffusion. The Al consumed during surface oxidation is related to the oxidation rate while the Al consumed from interdiffusion is dictated by the Al diffusivity in the alloy. The comparison in Table 4 of Co- and Ni-rich alloys without
any Re addition showed similar lifetimes at 1000 °C. The microstructure of these alloys consisted of the c and b phases, with the Co-rich alloys being more b rich. Since b is the main Al reservoir, the alloy with a higher b fraction should have a higher chemical life. However, the Co-rich alloys oxidized faster, similar to that reported by Pettit and Giggens [42] showing a factor of two higher Al2O3 growth rate on CoCrAl than on NiCrAl. Table 4 also shows that the Al diffusivity is faster in the Co-rich alloy, resulting in higher Al loss to the substrate (Table 3). A strong affinity between Al and Ni to form b-NiAl [43] is probably the cause for the lower Al diffusivity in the Ni-rich alloys. Therefore, although the Co-rich alloy had more starting b, it had a greater Al consumption rate due to higher oxidation and Al back diffusion rates, so the final lifetime in the Co- and Ni-rich alloys varied only by a few hundred hours. The Al content of the alloy is one of the most important variables since it is required for the formation of the protective Al2O3 scale; however, its amount in the alloy decreases with time due to depletion by oxidation and substrate interdiffusion. Increasing the alloy’s Al concentration resulted in a higher b fraction, thereby a longer coating life, as seen in Table 4. It is also apparent that Re additions increased the alloy’s b content as well. Furthermore, both the increase in Al and the addition of Re decreased the oxidation rates and the Al alloy diffusion rates. From literature review, it is anticipated that alloys with higher Al contents will produce purer Al2O3 during the early stage of oxidation [27] due to less growth of Co, Ni-containing oxides. Fig. 6 clearly shows that the mass gains at the end of the transient stages
Fig. 10. Calculated (solid lines) and experimentally determined (square symbols) Al concentrations in the CoRe-10 and NiRe-10 alloys after 500 h oxidation at 1000 °C. The vertical dashed line marks the thickness of the b-depletion layer, determined from the Al concentration above which the b phase becomes stable.
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Table 3 Chemical lifetime of coating alloys determined at 1000 °C. Lifetimes with oxidation alone and with Al back diffusion into a CMSX-10 substrate are presented. Alloy ID
CoRe-8 CoRe-10 CoRe-12 Co-10 NiRe-10 Ni-10
Coating life (h) Oxidation only
Oxidation and interdiffusion
% decrease in life with interdiffusion
9900 13,800 17,600 11,200 12,400 10,700
7800 11,700 15,500 8900 10,300 8600
21 15 12 21 17 20
Fig. 11. Calculated diffusion couple between CoRe-10 alloy and CMSX-10 substrate after 500 h at 1000 °C. Depletion zone-II represent the single phase region of the coating alloy.
were lower in alloys with higher Al. Alloys that develop lower quantities of spinels during the transient stage would produce a purer Al2O3 scale that grows slower than those that have large quantities of spinels incorporated within the scale. The same effects were also found for the Re-free and the Re-containing alloys. The fact shows that the steady-state oxidation rates of the Re-containing alloys could also be related to their lower transient-stage mass gains. Previous studies have reported significant decrease in the amount of spinels formed during the initial stage of oxidation of alloys containing Re and, consequently, lower steady-state rates as compared to the Re-free alloys [27]. Alloys with lower Al diffusivity are slightly less depleted of the b phase during oxidation and experience lower Al interdiffusion. Al diffusivity in the studied alloys decreased with additions of both Re and Al (Table 4). The decrease in DAl can be attributed to two causes. First, adding Re to the alloy substitutionally replaces the atoms in the c matrix with the larger Re. Secondly, Cr–Re particles
precipitated at the c/b phase boundaries in alloys with P 10% Al. These particles are r and a in the Co and Ni-rich alloys respectively. The r particles are a hard intermetallic TCP (tetragonal closed packed) phase, whereas a phase is a BCC (body centered cubic) solid solution. Al diffusivity in the intermetallic phases is much slower than in a solid solution like c and a [44]. Furthermore, because Al is insoluble in r (Table 2), it must either move around these particles or diffuse very slowly through them. Thus, r reduces Al diffusion in the Co-rich alloys. On the other hand, Al can diffuse easily through the solid solution a phase as it does in c, so there is little contribution of these particles on Al diffusivity in the Ni-rich alloy. Hence, as Table 4 suggests, the effects of Re on DAl are more pronounced in the Co-rich system. An increase of the alloy Al content also decreased the Al diffusivity in the alloy. However, the decrease is only substantial with the presence of precipitated r phase in alloys containing P 10% Al. This result again demonstrated the importance of alloy microstructure on diffusion. The role of the precipitated particles on Al diffusivity was demonstrated in Fig. 8. It is true that in the lower temperature range the thickness of the b-depletion layer will be smaller. Nevertheless, this thickness should increase with temperature at a constant rate unless the activation energy for DAl and/or for oxidation changes with temperature. The oxidation activation energies are known to remain constant [45] and the DAl data for a similar composition but with no Cr–Re precipitates showed that the activation energy actually remains constant in the temperature range of 900–1200 °C [12]. The kink in Fig. 8a shows that the increase in depletion thickness with increasing temperature is affected by a change in alloy microstructure. Since there are no reasons to believe the oxidation activation energies were affected by r or a dissolution, the effect on depletion thickness is considered to be an effect of diffusion activation energy. The different behaviors in Fig. 8a and b confirm the strong effect of the r phase on DAl in the Co-rich alloy, with no effect of a in the Ni-rich alloy. In a multicomponent system, like the Al2O3 forming Co-rich and Ni-rich alloys studied here, compositional change, no matter how insignificant, can lead to noticeable microstructural variations. Lower kp and DAl values together with higher b fraction resulted in higher chemical lives in alloys containing Re and higher Al. The final lifetimes depend on the coating composition and microstructure, operating temperature, coating thickness and the substrate chemistry. Fig. 12 compares the calculated lifetime of similar alloys of varying thicknesses deposited on a variety of substrates. The lifetimes calculated in the present work turned out to be very similar to the lifetimes reported in the literature [19,20]. Results presented in this figure also reflects the apparent life-increasing effect of Re on these alloys. Note that the lifetime calculated here represents the chemical life of the alloy and does not consider oxide film spallation. The lifetime should be lower if the
Table 4 Summary of the studied experimental and calculated aspects of bond coat oxidized at 1000 °C. Effect
a b
Alloy type
Microstructurea
Kp (g2 cm4 s1) 14
DAl (cm2 s1) 11
Lifetimeb(h)
Co vs. Ni
Co-rich alloy Ni-rich alloy
c + 0.47b c + 0.43b
30.6 10 18.6 1014
62.0 10 34.0 1011
8900 8600
Re addition in Co and Ni systems
Co-rich alloy Co alloy + Re Ni-rich alloy Ni alloy + Re
c + 0.47b c + 0.50b + 0.10r c + 0.43b c + 0.52b + 0.10a
30.6 1014 6.2 1014 18.6 1014 13.4 1014
62.0 1011 11.2 1011 34.0 1011 25.1 1011
8900 11,700 8600 10,300
Al content variation in Co-rich Re containing alloys
8% Al 10% Al 12% Al
c + 0.43b c + 0.50b + 0.10r c + 0.68b + 0.21r
15.6 1014 6.2 1014 4.0 1014
28.0 1011 11.2 1011 8.5 1011
7800 11,700 15,500
Phase fraction, in mole fraction, at 1000 °C. Calculated lifetime with oxidation and coating-substrate interdiffusion.
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References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]
Fig. 12. Lifetime of Co32Ni21Cr10Al3.5ReY alloy prepared in present work on CMSX-10 substrate. Lifetimes are compared with works of Krukovsky et al. [19,20].
[12] [13] [14] [15] [16] [17] [18] [19]
oxide spalls, e.g. during thermal cycling and the coating life would thus be limited by the thermomechanical stability of the TBC system [25,26].
[20]
5. Conclusions
[21] [22]
In the present work, the effects of Al and Re additions on the chemical lifetimes of several Co-rich and Ni-rich alloys were reported and analyzed under oxidizing conditions. Al depletion from the alloy, due to surface oxidation and coating-substrate interdiffusion, was modeled using DICTRA to determine every alloy’s final lifetime when the alloy was fully depleted of its b phase. The following conclusions can be drawn: 1. Without any Re addition, the Co-rich CoNiCrAlY alloy had a higher b content but lost more Al from oxidation and interdiffusion than the Ni-rich NiCoCrAlY alloy. Consequently, the chemical life of both alloys at 1000 °C varied by only a few hundred hours. 2. Alloy compositions containing Re provided significantly longer chemical lifetimes in both the Co- and the Ni-rich alloys. This improvement resulted from an increase in the alloy b content, a lowered steady-state Al2O3 growth rate and a decreased Al alloy diffusion rate. 3. Re addition is more effective in the Co- than the Ni-rich alloy. The difference in lifetime is mainly a result of the alloy phases, with the Co-rich alloy containing c, b, and r, and the Ni-rich alloy having c, b, and a. 4. The concentration of Al in an alloy has a profound effect on its lifetime. This is not only related to an increased b content but also to a lowered oxidation rate and a slower alloy Al diffusion rate.
Acknowledgments The authors are grateful for the financial support by National Natural Science Foundation of China (NSFC No. 51471094).
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Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Compositional effects on the high-temperature oxidation lifetime of MCrAlY type coating alloys S. Salam a, P.Y. Hou b, Y.-D. Zhang a, H.-F. Wang a, C. Zhang a, Z.-G. Yang a,⇑ a b
Key Laboratory of Advanced Materials, Ministry of Education, School of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA
a r t i c l e
i n f o
Article history: Received 17 January 2015 Accepted 9 March 2015 Available online 14 March 2015 Keywords: A. Alloy B. Modeling studies C. Oxidation
a b s t r a c t The lifetimes due to chemical failure during high-temperature oxidation of Co and Ni-rich MCrAlY type coating alloys were investigated. They were found to depend critically on the oxidation rate and Al diffusivity in alloy. Re and Al additions increased the life by providing higher b-phase contents and by lowering the Al diffusivity and oxidation rates. The effects of Re were more pronounced in the Co-rich than the Ni-rich alloys. The observed difference was a result of different alloy phases, with a microstructure of c, b, and r in the Co-rich alloys and c, b, and a in the Ni-rich alloys. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction Increase in the operating temperatures of modern gas turbines for improved efficiency have resulted in higher thermal–mechanical fatigue and oxidation of the turbines’ superalloy components [1,2]. To combat these problems, thermal barrier coatings (TBCs), consisting of a ceramic top coat (TC) and a metallic bond coat (BC), are usually applied to the substrates [2,3]. MCrAlY alloys, where M = Co, Ni or their combination, are one of the commonly used BC material. A thin surface layer of a slow-growing Al2O3 is developed between the TC and the BC at high temperatures by the preferential oxidation of Al in the MCrAlY [4–6]. This leads to the consumption of Al such that a subscale region is formed ahead of the Al2O3 layer, within which Al is depleted [7–9]. The consequent compositional changes in the alloy can give rise to microstructural changes where prolonged exposures may cause TBC failure [10,11]. TBCs applied in modern gas turbines, operating at a bond coat temperature of 900–1000 °C, require a lifetime of 25,000 h or more [12]. The TBC life is limited by a number of potential mechanisms including thermal fatigue failure, thermal aging failure, and Al depletion failure [12,13]. Out of these three, the Al depletion failure is considered as chemical breakdown, which occurs when the Al content in the alloy reaches a critically low value that can no longer sustain the continued growth of the protective scale. Consequently, less protective oxides, like fast-growing Ni(Co, Cr)spinels, are formed [14]. Experimental determination of the TBC ⇑ Corresponding author. Tel.: +86 1062783848; fax: +86 1062771160. E-mail address: [email protected] (Z.-G. Yang). http://dx.doi.org/10.1016/j.corsci.2015.03.011 0010-938X/Ó 2015 Elsevier Ltd. All rights reserved.
lifetime is not only tedious but also expensive. Instead, numerical analysis is preferred, which describes the alloy phase equilibria from calculated concentration changes in the alloy during hightemperature oxidation [15–18]. Often the time taken for an Al2O3 forming alloy to be completely depleted of its b phase is accepted as its chemical lifetime [19,20]. Recently, computer aided technologies for the modeling of phase diagrams and diffusion processes, like Thermo-Calc and DICTRA, have shown promising results to calculate the concentration changes underneath the scale during oxidation [21–23]. Microstructural changes arising from compositional variations could affect key bond coat properties, e.g. oxidation rate and diffusivity, which would affect the service life of the alloy. Oxidation rate of the bond coat alloy is clearly an important parameter, since the alloy with a higher rate would result in a more rapid consumption of the scale forming element. Diffusion in the alloy may also be important since it can deplete the alloy of essential constituents, such as Al, due to coating/substrate interdiffusion [11,24]. Existing experimental lifetime values are usually reported for thermal cycle failures [25,26], which are much lower than chemical lifetimes. Published chemical lifetime data vary substantially among different TBC systems [19,20]. These differences are attributed to parameters such as bond coat compositions, coating thicknesses, substrate compositions, operating temperatures, and deposition methods. Many researchers have attempted to improve the alloy’s oxidation resistance and lower the consumption of the scale forming element by improving the composition of the bond coat alloy [27–32]. However, the effects of microstructural phases independent of composition have received relatively little attention.
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In this study, CoNiCrAlReY alloys with different compositions were chosen. Prior works on these alloys had examined the effects of Al concentration and Re addition on their oxidation behavior [27], and had modeled the alloy phases resulted from these compositional changes [33]. The aim here is to investigate the influence of such compositional changes on the alloys’ chemical lifetime. Furthermore, in an effort to differentiate the effects of alloy phases from its composition, a number of Co- or Ni-rich MCrAlReY alloys were prepared, which provided alloys with different phase fractions of c + b with or without a or r precipitates. The chemical lifetimes of these alloys were calculated using Thermocalc and DICTRA. The lifetimes were then correlated with the effects of phases and compositions on the alloys’ oxidation and Al diffusion rates.
Fig. 1. Use of EPMA line scans to collect average concentration profiles over different positions along the diffusion direction in the alloy.
diffraction (XRD) (Rigaku/SmartLab) using Cu ka radiations at 40 kV.
2. Experimental method 2.1. Experimental work
2.2. Modeling procedure
Co and Ni-rich alloys were melted in a vacuum electric arc furnace (SK-II non-consumable vacuum arc furnace) from weighted percentage of pure elements in the form of high-purity (99.99%) metallic ingots obtained from Beijing Jiaming Platinum Industry of Non-Ferrous Metals Ltd. Co and Ni-rich alloys were melted to form button-shaped ingots of 100 diameter weighing 80 g each. Following melting, the ingots were homogenized in vacuum at 1150 °C for 24 h followed by the air cooling to room temperature. Actual compositions of the experimental alloys, determined with inductively coupled plasma mass spectrometer (ICP-MS) (Thermo Scientific™ iCAP™ Qc ICP-MS), are listed in Table 1. These alloys are given short names for simplicity, e.g. CoRe-8 means Co-rich alloy containing Re with 8% Al. Small disc-shaped samples (£10 mm 1 mm) were cut using an electrical discharge machine (EDM). Samples were mechanically ground and polished down to 0.5 lm grit using diamond paste. Finished samples were placed in alumina boats, containing 5 samples each, and inserted into a box furnace (Nabertherm LHT04/17) for oxidation in the air. Alloys were oxidized at 1000 °C for 100, 250, 500 and 1000 h. Additional samples were oxidized at 950, 1050, 1100, 1150, 1200 and 1250 °C for 100 h. For cross-sectional examination, oxidized samples were mounted in epoxy in a vertical orientation and then polished. The concentration profiles in subscale region were measured with an electron probe micro-analyser (EPMA) (Jeol JXA-8230) coupled with wavelength dispersive spectrometer (WDS). In order to accurately obtain the average concentration of different elements along their diffusion trajectory, a number of line scans were preformed parallel to the oxide/alloy interface as shown in Fig. 1. The length of each line was 100 lm, and the lines were spaced 2 lm apart. As seen from Fig. 1, these line scans pass through different phases; the elemental concentration at the location of each line was averaged over all the phases. Microstructures and phase compositions of the alloys were determined with a scanning electron microscope (SEM) (Jeol JSM-7001F). Alloy phases were identified with X-ray
Al consumption and its distribution in the alloys was modeled with DICTRA and Thermo-Calc using the thermodynamic and kinetic parameters obtained from the databases TTNI8 [34] and mob2 respectively. A continuous matrix (c phase) was considered containing b, and other precipitate phases dispersed within. The model assumes that diffusion takes place in the c matrix. The precipitate phases affect the diffusion process by their chemical composition and volume-fraction. Al was removed from the alloy with a flux proportional to the oxidation rate (kp) [35,36] using a hypothetical diffusion couple between an imaginary vacuum and the alloy. The vacuum acted as a sink for the Al leaving the alloy during oxidation. For a small time step Dt, Al in the alloy is removed through the diffusion couple, and the diffusion equations for the matrix phase are solved. New compositions are obtained at each point in the alloy assuming local equilibrium. For the next time step Dt, DICTRA solves the diffusion equations gain for the newer composition, etc. Using this simple modeling approach, the Al concentration and its distribution in the alloys was calculated as a function of oxidation time. Corresponding phase fractions were obtained from concentration changes within the alloy using calculated phase diagrams. Coating thickness is an important parameter while determining the total amount of b in the alloy. A typical bond coat thickness of 200 lm [19,20] was used to carry out the calculation. The total amount of the b phase in the alloy was continually calculated; the time taken for a complete b dissolution provides the chemical lifetime of the alloy.
Table 1 Composition of experimental alloys, determined by ICP-MS. Alloy ID
CoRe-8 CoRe-10 CoRe-12 Co-10 NiRe-10 Ni-10
Composition in wt.% Co
Ni
Cr
Al
Re
Y
Bal. Bal. Bal. Bal. 8.1 ± 0.4 8.2 ± 0.3
32.2 ± 0.2 31.9 ± 0.5 32.1 ± 0.3 31.9 ± 0.2 Bal. Bal.
21.3 ± 0.9 21.1 ± 0.7 20.9 ± 0.5 21.3 ± 0.3 20.9 ± 0.8 21.2 ± 0.1
8.1 ± 0.4 9.9 ± 0.6 11.8 ± 0.7 10.1 ± 0.3 9.9 ± 0.2 10.0 ± 0.2
3.4 ± 0.1 3.5 ± 0.3 3.4 ± 0.4 0 3.5 ± 0.1 0
0.13 0.15 0.12 0.11 0.12 0.13
3. Results 3.1. Microstructural characteristics Fig. 2 shows typical microstructures of the experimental alloys after they were heat-treated at 1000 °C for 500 h; corresponding XRD results are shown in Fig. 3. All alloys consisted mainly of the c and b phases, which appeared as light and dark areas respectively. c phase appeared as a continuous matrix phase. The microstructure of alloys without any Re addition (Co-10 and Ni-10) consisted of c and b phases only. Many white particles were observed at the c and b phase boundaries of the Re containing alloys. XRD analysis identified these particles as r phase in the Co-rich alloys and a phase in the Ni-rich alloys. No Ni–Y intermetallic phases were detected using XRD, because of the relatively low Y content in these alloys. Composition of several white particles (a or r) were analyzed with WDS, and their average concentrations are given in Table 2.
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Fig. 2. Backscattered electron images (BEI) of experimental alloys after heat-treatment at 1000 °C for 500 h; (a) CoRe-8, (b) CoRe-10, (c) CoRe-12, (d) Co-10, (e) NiRe-10, (f) Ni-10.
Fig. 3. XRD patterns of heat-treated Co-rich and Ni-rich alloys with and without Re. Alloys were heat-treated at 1000 °C for 500 h.
Both types of particles were enriched with Cr and Re, whose concentrations were higher in a than in r. The r particles also contained a significant amount of Co; no Y was detected in r and the Al concentration in them was very low. The a phase, on the other hand, contained more Al and a notable amount of Y.
The influence of heat-treatment temperature on alloy phases is demonstrated in Fig. 4. The amount of b and a/r phases decreased with increasing temperature, and the a/r phases eventually dissolved, due to increased solubility of Cr and Re in c at higher temperatures [37]. A binary c + b microstructure was observed
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Table 2 WDS analysis showing the average composition of white particles in Fig. 2. White particles
r phase a phase
Composition in wt.% Co
Ni
Cr
Al
Re
Y
26.8 ± 0.9 2.5 ± 1.2
7.2 ± 0.4 5.9 ± 2.6
41.4 ± 0.3 56.6 ± 1.0
0.3 ± 0.1 1.7 ± 0.2
24.3 ± 0.6 30.4 ± 1.2
0 2.96 ± 0.4
for the CoRe-10 and NiRe-10 alloys heat-treated at P 1150 °C and 1200 °C respectively. The relative amount of different phases in each alloy was estimated using image analysis. A detailed method of obtaining the weighted percentage of phases from their measured surface fraction is given in an earlier paper [33]. The calculated phase amounts are shown as a function of heat-treatment temperature in Fig. 5. Co and Ni-rich alloys had similar phase content at all tested temperatures that consisted of c and b only. All six alloys showed decreasing fractions of b and r or a with increasing temperature. For a given temperature, higher amounts of b were present in alloys with more Al (compare CoRe-8, CoRe-10, and CoRe-12). The r phase precipitated in alloys with Al P 10% and increased with increasing Al content in the alloy. Re addition not only resulted in the precipitation of r or a, it also increased the alloy’s b content. This effect was slightly more pronounced in the Ni-rich than in Corich alloy. 3.2. Oxidation and diffusion All investigated alloys underwent uniform oxidation. XRD analysis through each alloy’s surface confirmed that the oxide was mainly a-Al2O3 with small amounts of Co,Ni spinels formed during the initial stages of oxidation. The oxidation rate of the alloys during the transient stage was significantly higher than in the steady-state (>10 h). A few examples of sample mass gains as a function of oxidation time are shown in Fig. 6a. A quadratic curve with substantial mass gain during the first few hours can be seen. Similar two stage oxidation behavior has been extensively reported in the literature for Al2O3 forming alloys [38,39] where initially fast
growing (Co, Ni)-spinels and transition Al2O3 form on alloy surface. A continuous a-Al2O3 does not form until at least few hours have elapsed [40,41]. Consequently, in the transient stage the slope of mass gain curves is much steeper than those seen in the steadystate. To approximate the rate constants for a quantitative comparison, the relatively parabolic section of the curve was fitted to the parabolic rate equation given by [40]:
Dm 2 ¼ kp t þ C A
ð1Þ
where Dm is the mass gain, A is the specimen’s surface area, kp is the rate constant and t is the oxidation time. Fig. 6b plots the square of the mass gains vs. the oxidation time in the second stage. Parabolic time dependency of the a-Al2O3 growth is seen for all studied alloys. The first data point of each curve in Fig. 6b represents the end of the transient stage oxidation. Approximate parabolic rate constants (kp) for the steady-state were obtained from the slopes of the lines in the parabolic plot. These kp values are indicated in the figure caption, and they are in good agreement with those reported for a-Al2O3 growth in similar alloys [42]. Fig. 6b shows that the alloys with more Al oxidized slower. The Co-rich alloys oxidized at a slightly higher rate than the Ni-rich alloys. Re addition, decreased the steady-state rate in both types of alloys, but the effect was stronger in the Co- than in Ni-rich alloys. After oxidation, polished cross-sections were examined using SEM. A dense and uniform Al2O3 scale was seen on the alloy’s surface in all the test samples. Beneath the scale, a depth containing only the c phase, identified as a b-depletion layer, was observed in all alloys. An example is shown in Fig. 7 with the CoRe-10 alloy. It is seen that the depth of the b-depletion layer
Fig. 4. Microstructure of CoRe-10 and NiRe-10 alloys after heat-treatments at different temperatures; (a–c) CoRe-10 alloy at 1100 °C, 1150 °C and 1200 °C, respectively; (d–f) NiRe-10 alloy at 1100 °C, 1150 °C and 1200 °C, respectively.
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Fig. 5. Phase content of experimental alloys, in wt.%, as a function of heat treatment temperature for (a) b, (b) c, (c) r and (d) a. Solid lines represent the Co-rich alloys and dotted lines represent the Ni-rich alloys.
increased with increasing temperature and oxidation time. Slight accumulation of the Cr–Re particles were observed at the scale/alloy interface of the alloy oxidized at 1000 °C; this behavior has been documented and explained in an earlier paper [35]. The thickness of the b-depletion layer after 100 h oxidation in CoRe-10 and NiRe-10 alloys was measured in the temperature range between 1000 and 1250 °C. These results are presented in Fig. 8, where the layer thicknesses are shown as a function of 1/T. Vertically dashed lines are used to indicate the boundary between two microstructure types, i.e. c + b at higher temperatures and c + b + r or c + b + a at lower temperatures. As the depletion thickness increases with increasing temperature, linear behavior was expected in the Arrhenius plot. This was indeed the case for the NiRe-10 alloy, but for the CoRe-10 alloy, two distinct slopes were present, suggesting that the depletion rates in the two regions were different. Within the temperature range of 1000 and 1100 °C, the microstructure consisted of c + b + r phases and the Co-rich alloy was depleted at a lower rate than in the temperature range of 1100 and 1250 °C where no r was present. The depletion rate in the Ni-rich alloy remained the same throughout the tested temperature range. The Al diffusivity in the c matrix was estimated using the analytical solution proposed by Whittle through the equation [14]:
rffiffiffiffiffiffiffiffi 1 pkc X nL ðn þ 1ÞL Ni ¼ No erfc pffiffiffiffiffiffi þ erfc pffiffiffiffiffiffi 2D n¼0 2 Dt 2 Dt
ð2Þ
where Ni is the experimentally determined interface Al concentration, No the Al concentration in the alloy, kc the oxidation rate constant in (cm2 s1), L the sample thickness, t the oxidation time and D
is the diffusion coefficient. To keep the units consistent, the measured kp was converted to kc using the equation:
kc ¼ 2
Mx VZ x
2
kp
ð3Þ
where Mx is the atomic mass of oxygen, V the volume of the oxide scale and Zx the valency of oxygen. The calculated values of DAl in the gamma matrix for each of the studied alloys are presented in Fig. 9. Data shows that Al diffused faster in the Co-rich than the Ni-rich alloy. Increasing the alloy Al content from CoRe-8 to CoRe12 decreased the Al diffusivity. Re addition also decreased the Al diffusivity; the effect was considerably stronger in the Co-rich alloy than in Ni-rich alloy.
3.3. Lifetime prediction Al concentration profiles in the alloy after different oxidation times were calculated using the method explained in the modeling section. To verify the accuracy of the modeling method, profiles for each alloy were calculated for 250 and 500 h oxidations at 1000 °C and compared with experimental values determined using EPMA; good matches were observed between the two. An example is shown in Fig. 10, where Al concentrations (solid lines from modeling and data points from EPMA measurement) in CoRe-10 and NiRe-10 alloys are plotted against the distance from the scale/alloy interface into the bulk alloy. The width of the b-depletion layer is marked as a vertical dash line, whose location was determined at the Al concentration above which b becomes stable (based on the phase diagrams calculated using Thermo-Calc). The thicknesses
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Fig. 6. Oxidation kinetics of the experimental alloys at 1000 °C; (a) mass gain as a function of oxidation time; (b) parabolic plot of (a). The slope of each curves in (b) represents the steady-state kp of the alloy. Calculated kp values in g2 cm4 s1 are shown in the legend.
of these calculated b-depletion layers were in good agreement with those observed from the sample cross-sections. The b phase in the alloy is consumed with increasing oxidation time until its content reaches zero. Lifetimes predicted by the complete consumption of each alloy’s b phase are presented in Table 3.
These modeled results show that alloys with larger b reservoirs (more Al) have longer lives; Co- and Ni-rich alloys have similar lifetimes that increase with Re addition; Re in the Co- and Ni-rich alloys increases their lifetimes by 2600 h and 1700 h respectively. Table 3 also includes the effect of coating-substrate interdiffusion on lifetime. When Al in the bond coat diffuses into the substrate, less Al is available to continue the growth of the protective Al2O3 scale. Diffusion couples between the tested alloys and a third-generation superalloy substrate (CMSX-10 with about 6% Al) were simulated with DICTRA to determine the extent of Al interdiffusion between these two. An example of the resulting diffusion profile, between CoRe-10 and CMSX-10 after 500 h at 1000 °C, is shown in Fig. 11. The vertical dash line in the figure represents the original interface between the coating alloy and the substrate. As seen from Fig. 11, the Al back-diffuses into the substrate and the Al concentration in the coating decreases. A subsequent Al depleted single c phase zone (the marked depletion layer-II on Fig. 11) is formed, and its depth increases with time. This profile indicates that Al in the single crystal substrate with no grain boundaries diffuses at a slower rate [12], while the polycrystalline coating alloys with grain boundary diffusion allowed faster diffusion in the coating. Calculated width of interdiffusion zone in the coating did not extend more than 50 lm. For the lifetime modeling, the width of this depletion layer-II at any given time was calculated and then subtracted from the coating’s original thickness. This correction undoubtedly reduces the total amount of the b content in the alloy and consequently reduces the lifetime. The corrected lifetimes and their % decrease from the oxidation-only condition without any diffusion into a substrate are shown in Table 3. The values in the table show that coating lives decrease by 12–21% due to Al back-diffusion into the substrate. The extent of this back diffusion decreases with increasing Al content in the alloy and with Re addition.
4. Discussion Minor changes to the alloy’s composition affect the bond coat properties that dictate the chemical life of the alloy. Results from this study are summarized in Table 4, which show the effects of
Fig. 7. SEM cross-sectional images of CoRe-10 alloy oxidized at (a) 1000 °C for 100 h, (b) 1000 °C for 500 h, (c) 1100 °C for 100 h and (d) 1200 °C for 100 h. A b-depletion layer can be seen beneath the Al2O3 scale under all conditions.
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Fig. 8. Temperature dependence of the b-depletion thickness in Re-containing alloys oxidized for 100 h; (a) CoRe-10, (b) NiRe-10.
Fig. 9. Al diffusivity in the experimental alloys calculated at 1000 °C.
alloy chemistry and its resulting microstructure on the oxidation rate, the Al alloy diffusion rate and the chemical lifetime of tested Co-rich and Ni-rich coating alloys. The lifetime of an alloy can be increased by adding Al or by modifying its composition to reduce the Al consumption during subsequent oxidation and interdiffusion. The Al consumed during surface oxidation is related to the oxidation rate while the Al consumed from interdiffusion is dictated by the Al diffusivity in the alloy. The comparison in Table 4 of Co- and Ni-rich alloys without
any Re addition showed similar lifetimes at 1000 °C. The microstructure of these alloys consisted of the c and b phases, with the Co-rich alloys being more b rich. Since b is the main Al reservoir, the alloy with a higher b fraction should have a higher chemical life. However, the Co-rich alloys oxidized faster, similar to that reported by Pettit and Giggens [42] showing a factor of two higher Al2O3 growth rate on CoCrAl than on NiCrAl. Table 4 also shows that the Al diffusivity is faster in the Co-rich alloy, resulting in higher Al loss to the substrate (Table 3). A strong affinity between Al and Ni to form b-NiAl [43] is probably the cause for the lower Al diffusivity in the Ni-rich alloys. Therefore, although the Co-rich alloy had more starting b, it had a greater Al consumption rate due to higher oxidation and Al back diffusion rates, so the final lifetime in the Co- and Ni-rich alloys varied only by a few hundred hours. The Al content of the alloy is one of the most important variables since it is required for the formation of the protective Al2O3 scale; however, its amount in the alloy decreases with time due to depletion by oxidation and substrate interdiffusion. Increasing the alloy’s Al concentration resulted in a higher b fraction, thereby a longer coating life, as seen in Table 4. It is also apparent that Re additions increased the alloy’s b content as well. Furthermore, both the increase in Al and the addition of Re decreased the oxidation rates and the Al alloy diffusion rates. From literature review, it is anticipated that alloys with higher Al contents will produce purer Al2O3 during the early stage of oxidation [27] due to less growth of Co, Ni-containing oxides. Fig. 6 clearly shows that the mass gains at the end of the transient stages
Fig. 10. Calculated (solid lines) and experimentally determined (square symbols) Al concentrations in the CoRe-10 and NiRe-10 alloys after 500 h oxidation at 1000 °C. The vertical dashed line marks the thickness of the b-depletion layer, determined from the Al concentration above which the b phase becomes stable.
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Table 3 Chemical lifetime of coating alloys determined at 1000 °C. Lifetimes with oxidation alone and with Al back diffusion into a CMSX-10 substrate are presented. Alloy ID
CoRe-8 CoRe-10 CoRe-12 Co-10 NiRe-10 Ni-10
Coating life (h) Oxidation only
Oxidation and interdiffusion
% decrease in life with interdiffusion
9900 13,800 17,600 11,200 12,400 10,700
7800 11,700 15,500 8900 10,300 8600
21 15 12 21 17 20
Fig. 11. Calculated diffusion couple between CoRe-10 alloy and CMSX-10 substrate after 500 h at 1000 °C. Depletion zone-II represent the single phase region of the coating alloy.
were lower in alloys with higher Al. Alloys that develop lower quantities of spinels during the transient stage would produce a purer Al2O3 scale that grows slower than those that have large quantities of spinels incorporated within the scale. The same effects were also found for the Re-free and the Re-containing alloys. The fact shows that the steady-state oxidation rates of the Re-containing alloys could also be related to their lower transient-stage mass gains. Previous studies have reported significant decrease in the amount of spinels formed during the initial stage of oxidation of alloys containing Re and, consequently, lower steady-state rates as compared to the Re-free alloys [27]. Alloys with lower Al diffusivity are slightly less depleted of the b phase during oxidation and experience lower Al interdiffusion. Al diffusivity in the studied alloys decreased with additions of both Re and Al (Table 4). The decrease in DAl can be attributed to two causes. First, adding Re to the alloy substitutionally replaces the atoms in the c matrix with the larger Re. Secondly, Cr–Re particles
precipitated at the c/b phase boundaries in alloys with P 10% Al. These particles are r and a in the Co and Ni-rich alloys respectively. The r particles are a hard intermetallic TCP (tetragonal closed packed) phase, whereas a phase is a BCC (body centered cubic) solid solution. Al diffusivity in the intermetallic phases is much slower than in a solid solution like c and a [44]. Furthermore, because Al is insoluble in r (Table 2), it must either move around these particles or diffuse very slowly through them. Thus, r reduces Al diffusion in the Co-rich alloys. On the other hand, Al can diffuse easily through the solid solution a phase as it does in c, so there is little contribution of these particles on Al diffusivity in the Ni-rich alloy. Hence, as Table 4 suggests, the effects of Re on DAl are more pronounced in the Co-rich system. An increase of the alloy Al content also decreased the Al diffusivity in the alloy. However, the decrease is only substantial with the presence of precipitated r phase in alloys containing P 10% Al. This result again demonstrated the importance of alloy microstructure on diffusion. The role of the precipitated particles on Al diffusivity was demonstrated in Fig. 8. It is true that in the lower temperature range the thickness of the b-depletion layer will be smaller. Nevertheless, this thickness should increase with temperature at a constant rate unless the activation energy for DAl and/or for oxidation changes with temperature. The oxidation activation energies are known to remain constant [45] and the DAl data for a similar composition but with no Cr–Re precipitates showed that the activation energy actually remains constant in the temperature range of 900–1200 °C [12]. The kink in Fig. 8a shows that the increase in depletion thickness with increasing temperature is affected by a change in alloy microstructure. Since there are no reasons to believe the oxidation activation energies were affected by r or a dissolution, the effect on depletion thickness is considered to be an effect of diffusion activation energy. The different behaviors in Fig. 8a and b confirm the strong effect of the r phase on DAl in the Co-rich alloy, with no effect of a in the Ni-rich alloy. In a multicomponent system, like the Al2O3 forming Co-rich and Ni-rich alloys studied here, compositional change, no matter how insignificant, can lead to noticeable microstructural variations. Lower kp and DAl values together with higher b fraction resulted in higher chemical lives in alloys containing Re and higher Al. The final lifetimes depend on the coating composition and microstructure, operating temperature, coating thickness and the substrate chemistry. Fig. 12 compares the calculated lifetime of similar alloys of varying thicknesses deposited on a variety of substrates. The lifetimes calculated in the present work turned out to be very similar to the lifetimes reported in the literature [19,20]. Results presented in this figure also reflects the apparent life-increasing effect of Re on these alloys. Note that the lifetime calculated here represents the chemical life of the alloy and does not consider oxide film spallation. The lifetime should be lower if the
Table 4 Summary of the studied experimental and calculated aspects of bond coat oxidized at 1000 °C. Effect
a b
Alloy type
Microstructurea
Kp (g2 cm4 s1) 14
DAl (cm2 s1) 11
Lifetimeb(h)
Co vs. Ni
Co-rich alloy Ni-rich alloy
c + 0.47b c + 0.43b
30.6 10 18.6 1014
62.0 10 34.0 1011
8900 8600
Re addition in Co and Ni systems
Co-rich alloy Co alloy + Re Ni-rich alloy Ni alloy + Re
c + 0.47b c + 0.50b + 0.10r c + 0.43b c + 0.52b + 0.10a
30.6 1014 6.2 1014 18.6 1014 13.4 1014
62.0 1011 11.2 1011 34.0 1011 25.1 1011
8900 11,700 8600 10,300
Al content variation in Co-rich Re containing alloys
8% Al 10% Al 12% Al
c + 0.43b c + 0.50b + 0.10r c + 0.68b + 0.21r
15.6 1014 6.2 1014 4.0 1014
28.0 1011 11.2 1011 8.5 1011
7800 11,700 15,500
Phase fraction, in mole fraction, at 1000 °C. Calculated lifetime with oxidation and coating-substrate interdiffusion.
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Fig. 12. Lifetime of Co32Ni21Cr10Al3.5ReY alloy prepared in present work on CMSX-10 substrate. Lifetimes are compared with works of Krukovsky et al. [19,20].
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[20]
5. Conclusions
[21] [22]
In the present work, the effects of Al and Re additions on the chemical lifetimes of several Co-rich and Ni-rich alloys were reported and analyzed under oxidizing conditions. Al depletion from the alloy, due to surface oxidation and coating-substrate interdiffusion, was modeled using DICTRA to determine every alloy’s final lifetime when the alloy was fully depleted of its b phase. The following conclusions can be drawn: 1. Without any Re addition, the Co-rich CoNiCrAlY alloy had a higher b content but lost more Al from oxidation and interdiffusion than the Ni-rich NiCoCrAlY alloy. Consequently, the chemical life of both alloys at 1000 °C varied by only a few hundred hours. 2. Alloy compositions containing Re provided significantly longer chemical lifetimes in both the Co- and the Ni-rich alloys. This improvement resulted from an increase in the alloy b content, a lowered steady-state Al2O3 growth rate and a decreased Al alloy diffusion rate. 3. Re addition is more effective in the Co- than the Ni-rich alloy. The difference in lifetime is mainly a result of the alloy phases, with the Co-rich alloy containing c, b, and r, and the Ni-rich alloy having c, b, and a. 4. The concentration of Al in an alloy has a profound effect on its lifetime. This is not only related to an increased b content but also to a lowered oxidation rate and a slower alloy Al diffusion rate.
Acknowledgments The authors are grateful for the financial support by National Natural Science Foundation of China (NSFC No. 51471094).
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