Low cycle fatigue behavior of Platinum-Aluminide coated Rene^®80 near and above the DBTT

Mohammad Mehdi Barjesteh^a,*, Seyed Mehdi Abbasi^a, Karim Zangeneh Madar^a, Kourosh Shirvani^b

^aFaculty of Material and Manufacturing Technologies, Malek Ashtar University of Technology (MUT), Babaie street, Lavizan, 15875-1774, Tehran, Iran

^bDepartment of Advanced Materials and New Energies, Iranian Research organization for Science and Technology (IROST), 33535-111, Tehran, Iran

(*Corresponding author: mmbarjesteh@yahoo.com)

1. INTRODUCTIONTOP

Rene^®80 (General Electric trademark) is a cast nickel-based superalloy used for manufacturing of first stage turbine blades of jet engines. This superalloy is generally used at the temperature range of 760–982 °C (Safari and Nategh, 2006). Protective coatings are deposited on the surface of airfoil turbine blades, in order to enhance the corrosion and oxidation resistance of the base alloys at high temperatures. It has been known that the formation of a protective alumina scale based on intermetallic β-NiAl compound, improves the oxidation resistance of the base alloys (Rahmani and Nategh, 2008a). More improvement in cyclic oxidation and hot corrosion resistance of superalloys has been reported by controlling the addition of precious metals such as platinum to β-NiAl compound (Krishna et al., 1998). Although the presence of Pt in aluminide coatings causes an improvement in their cyclic oxidation resistance at high temperatures, a minimum content of Pt (corresponding to a 6 mm thick initial Pt layer) in the coating is necessary in order to realize its beneficial effects fully (Krishna et al., 1998). To produce the platinum-aluminide (Pt-Al) coating, an initial layer of platinum was applied on the considered surface and after that, aluminum is diffused into it. In the diffusion coatings, the differences in the composition of the coating and substrate, as a result of coating and the substrate element combinations, have a significant effect on the mechanical properties of the alloy. (Ni, Pt)Al and PtAl₂ phases which are formed during Pt-Al coating process, are inherently brittle and during service, are prone to cracking at low temperature. The DBTT (Ductile to Brittle Transition Temperature) of aluminide bond coats by micro-tensile test method was evaluated in a study Alam et al. (2010), They found that the Pt-Al coating has a higher DBTT than a plain aluminide coating. The DBTT values for various Pt-Al and plain aluminide coatings have been reported in the range of 850–1000 and 700–800 °C, respectively. The contents of Al and Pt in Pt-Al coatings have a significant effect on DBTT. It means that coatings with higher surface aluminum and platinum levels exhibited a higher DBTT. The formation of a continuous layer of PtAl₂ on the surface of Pt-Al coatings, is a reason for this phenomenon (Vogel et al., 1987). Depending on working temperature, below or above DBTT, negative or positive effects of the coating on mechanical properties of the base alloys are reported. The DBTT is the temperature at which the slope of the strain-temperature curve changes significantly. This temperature is affected by coating process, phase distribution, composition, heat treatment and microstructure. Diffusion coatings exhibit low ductility below the DBTT, whereas above this temperature ductility increases (Bose, 2007). Decreased low cycle fatigue (LCF) properties of Rene^®80 alloy following the application of CODEP-B, a plain aluminide coating, at the total strain range 1.2%, and an increase in this property at total strains below 0.8% at 871 °C have been reported by Rahmani and Nategh (2008b), while applying of Pt-Al coating on this superalloy reduced low cycle fatigue property at the same total strain and temperature conditions (Barjesteh et al., 2019a).

Since fatigue failures are in most cases initiated at the surface, in the absence of internal defects and porosity, a coating may potentially affect fatigue strength of coated components, therefore, to study the influence of coating on the fatigue of superalloy has important significance. Nowadays, designers of jet engines, increase the turbine inlet temperature (TIT) with the aim of improving the performance of jet engines. As a result, the fatigue behavior of coated rotary components, such as turbine blades, has been affected by temperature variations. In the present study, the effect of temperature (near and above the DBTT) on low cycle fatigue of a high activity Pt-Al coated Rene^®80 was evaluated.

2. MATERIALS AND METHODSTOP

The cast nickel base superalloy Rene^®80 (nominal composition: 0.16 C, 13.81 Cr, 9.69 Co, 4.23 Mo, 4.02 W, 3.02 Al, 4.87 Ti, 0.12 Fe, 0.05 Zr, 0.05 V, 0.03 Mn, 0.02 B, 0.02 Si, balance Ni, in wt.%, as determined by quantometer apparatus) was used as the substrate in this study. According to Fig. 1, fatigue specimens were manufactured by machining cylindrical bars of 24 mm in diameter and 120 mm in length obtained from investment casting, as per ASTM E606 (Cameron et al., 1996). Then, to make sure there are no surficial or internal defects, all specimens were inspected by the Fluorescent Penetration Inspection (FPI) and X-ray test method.

The next step involved a solution heat treatment at 1205 °C for 2 h and first aging treatment at 1095 °C for 4 h (Rush, 2012). After grit blasting and cleaning the specimens with acetone, an intermediate layer of nickel with a thickness of 1–2 µm was applied on the gauge area of the specimens for decreasing the negative effect of the chromium (available on the composition of the alloy) on the lack of adhesive property of platinum. Then, the gauge area of the specimens was coated with a layer of Pt using an electroplating method. Platinum plating was carried out using an electrolyte solution containing 14–18 ml of type P salt (di-nitro di-amino platinum), 70–90 g·L^-1 calcium carbonate (Na₂CO₃), 40–70 g·L^-1 sodium acetate (NaCH₃COO), and 1L of distilled water at 90 °C under a current density of 0.2-0.4 A·dm^-2 and electrolyte pH of about 10.5 (Rashidghamat and Shirvani, 2009). In order to achieve a platinum layer with a thickness of 6 µm, the time of 360 min was considered for the plating process. The platinum layer was subjected to heat treatment at 1050 °C for 2 h to enhance cohesion and distribution of the platinum into the substrate, followed by cooling the specimens in a furnace at 400 °C and then air-cooled (Pedraza et al., 2006). The aluminizing process was performed under Low Temperature-High Activity (LTHA), at 750 °C for 4 h followed by post aluminizing for 2 h at 1050 °C, via powder cementation. Composition of the cementation powder used for LTHA condition was selected as 2NH₄Cl (Activator) -12Al (Main Agent) -86Al₂O₃ (Filler). After the formation of the Pt-Al coating on the surface, an aging treatment was performed at 845 °C for 16 h (Rush, 2012).

Microstructural characterization and fractography studies were conducted using a scanning electron microscope (SEM) Model Zeiss Supra 55 equipped with energy dispersive spectroscopy (EDS) Oxford Model, both prior to performing fatigue test (to ensure the quality of the coating) and after the test (to investigate fracture surfaces) according to ASTM E3 (Brown and Donmez, 2016), ASTM E407 (Walker and Tarn, 1991), and ASTM E112-96 (Vander Voort, 2004) standards. X-ray diffraction analysis was performed using an Inel Equinox 6000 with X’Pert High Score Plus v2.0, Cu Kα1 with Graphite monochromator, 2θ = 16° to 93°, to determine distributions of different phases across the coating thickness. Micro-hardness testing was carried out normal to the alignment of the Pt-Al coating and the substrate according to ASTM E384 (Grote and Antonsson, 2010) using an automatic Akashi micro-hardness tester equipped with the Clemex software through applying 50 gf. Tensile tests were performed on coated and uncoated specimens at 982 °C using a Universal ATM apparatus equipped with an electrical furnace that could apply temperatures up to 1000 °C ± 1. The tensile tests were performed with a constant strain rate of 1.6×10^-4 s^-1, in accordance with ASTM E21 standard procedure (Kuhn and Medlin, 2000). Finally, low cycle fatigue tests were performed on the coated and uncoated specimens at 871 °C and 982 °C according to ASTM E606 standard test method (Cameron et al., 1996), ΔƐ_t = 0.8%, strain rate of 2×10^-3 s^-1, tensile-tensile triangular wave, and R = ε_min/ε_max =0. A servo-hydraulic Instron fatigue test machine equipped with an electrical furnace (1000 °C ± 1) and a high temperature extensometer used to conduct high temperature fatigue tests in strain control mode. It should be mentioned that prior to testing at elevated temperatures, the specimens were allowed to come to thermal equilibrium for 20 min at zero load.

3. RESULTS AND DISCUSSIONTOP

3.1. Microstructural characterizationTOP

The microstructures of the substrate (Rene^®80) and the coating are shown in Fig. 2. No difference was seen between the microstructure of coated and uncoated substrate.

The microstructure of the Rene^®80 Fig. 2a consisted of a γ-phase matrix, primary, and secondary γ´. Studies have indicated that there are no changes in the morphology, size and distribution of the γ´ (Ni₃(Al,Ti)) phase (strengthening precipitates) in both coated and uncoated substrates. Hardness of the coated and uncoated substrates was measured as 404 and 408 HV, respectively. The average grain size of the coated and uncoated substrate was found to be 0.4–0.6 mm (measured on a Clemex image analyzer v3.5) according to the ASTM E-112-96 standard (Vander Voort, 2004). Figure 2b shows the cross-sectional microstructure of the coating used in the present study (6 μm Pt/LTHA). The coating exhibited the typical three-layer structure. As shown in this figure, the outer layer contained ξ-PtAl₂ (bright contrast) and β-(Ni, Pt)Al. The presence of ξ-PtAl₂ and β-(Ni, Pt)Al phases, in this layer, was also confirmed from XRD analysis (Fig. 3). The intermediate layer constituted single-phase β-(Ni,Pt)Al and final layer named inter-diffusion zone (IDZ, the coating-substrate interface). The thicknesses of the outer layer, the intermediate layer and IDZ were measured 44 ±4, 92 ±7 and 4 ±1 μm, respectively. Hardness in the distance of 15 μm from the surface of the coating was measured 1011 HV, while this value was measured 728 HV in the depth of 125 μm from the coating surface. This high value of hardness shows the brittle nature of Pt-Al coatings.

The concentration plots for Pt, Al and Ni in the coating are shown in Fig. 4. Results of EDS line scan analysis are indicative of internal diffusion of platinum and aluminum and external diffusion of nickel. As shown in this figure, the surface Pt concentration is about 40 wt.%. Based on the report by Pedraza et al. (2006), at least 28 wt.% of platinum in the surface of Pt-Al coating is necessary to constitute a bi-phase structure.

3.2. Low cycle fatigue behaviorTOP

3.2.1. Monotonic stress-strain behavior (Tensile Properties)TOP

The tensile test of superalloy coated and uncoated Rene^®80 results and discussions obtained at 871 °C have been reported by us previously (Barjesteh et al., 2019b). The results are shown in Table 1. According to the results, the tensile properties of coated specimens are lower than uncoated specimens. In the present study, tensile test is evaluated on coated and uncoated specimens at 982 °C. The results of the test for this condition are reported in Table 2. It should be mentioned that tensile test at this temperature was performed on two uncoated specimens (after full heat treatment, according to GE-C50TF28 (AMS 5403) (Rush, 2012) and two coated (6 μm Pt/LTHA) specimens. As it can be observed obvious from the Tables 1 and 2, by increasing the testing temperature from 871 °C to 982 °C, in both uncoated and coated specimens, ultimate tensile strength (UTS) and yield strength (YS) have been rapidly decreased while ductility (elongation (El%) and reduction of area (RA%)) of uncoated and coated specimen increased. For example, coated specimen showed 55.4% lower YS at 982 °C compared to 871 °C while the elongation of coated specimen at 982 °C was about 64.2% higher than this property at 871 °C. The strength of Pt-Al coatings is lower than that of the substrate alloy on which they are deposited (Alam et al., 2011). For example, yield and ultimate strengths of Pt–Al coatings were equal to 200 and 300 MPa at a temperature range of 800–900 °C, respectively, which was lower than the yield strength (595 MPa) and ultimate strength (706 MPa) of the Rene^®80 alloy (Table 1) at 871 °C. Different fracture stress will be made with this huge difference between the coating and substrate strengths that give rise to the formation of some cracks on the coating during the usage of tensile load. Therefore, it can be expected that the presence of the coating would cause some degree of weakening of the base alloy.

Fractographic studies of the uncoated and coated (6 μm Pt/LTHA) tensile specimens at 871 °C in the previous investigation (Barjesteh et al., 2019b) exhibited that the fracture mechanism was a mixed mode of ductile and brittle fracture. The results of the previous study also showed that most of the cracks nucleate from grain boundaries and propagate along them in the uncoated specimens, while cracks initiated from the coating surface in the coated specimens. Figure 5 shows the SEM images of uncoated and coated specimens failed during a tensile test at 982 °C. As it can be observed from Fig. 5a, at higher temperature, not only grain boundaries and thick oxide layer are the suitable locations for crack nucleation, in the uncoated specimens, but also internal voids are other appropriate sites. With increase in test temperature, more dimples with the large size could be observed in the uncoated specimens (Fig. 5b). According to Fig. 5c, profuse large dimple formation also occurred in the coated specimens, at 982 °C. SEM image of the fracture surface (longitudinal section) at 982 °C of coated specimens is provided in Fig. 5d.

Extensive void formation and ductile failure were seen across the coating thickness. The fracture surface features of the superalloy substrate at high temperatures were similar to those of uncoated specimens. To study the crack behavior in coated specimens, the longitudinal section of these specimens after the tensile test was examined. According to Fig. 5e, the crack remained in the coating thickness range and no penetration of cracks into the substrate has been observed. Other studies (Parlikar et al., 2017) showed a decrease in the extent of crack penetration with the increase in test temperature. However, it is known that these cracks decrease the load bearing cross-section area of the substrate and cause a decrease in the strength of the substrate alloy. As mentioned previously, large dimples have been observed in the cross section of coated specimens, this was because the test temperature (982 °C was well above the DBTT of the coating. Comparison of obtained results with the previous investigation (Barjesteh et al., 2019b) showed that the extent of the ductile fracture region was much more at 982 °C.

3.2.2. Cyclic stress-strain behavior of uncoated specimens TOP

Results of fatigue test on the uncoated Rene^®80 specimens at 871 °C and total strain range of 0.8% have been reported by us in the previous study (Barjesteh et al., 2019a). The results of the previous study are shown in Table 3. Low cycle fatigue test of uncoated Rene^®80 at 982 °C, also performed in the present study and the results are shown in Table 4, where ΔƐ_t is total strain amplitude, σ_max is maximum cyclic stress, σ_min is minimum cyclic stress, N_i is the number of cycles by which the first drop in the alloy properties is observed (crack nucleation), and N_f is the number of cycles to fracture.

Table 3. Results of low-cycle fatigue test on uncoated Rene^®80 specimens (ΔƐ_t (%) =0.8, R =0, = 2 × 10^-3 s^-1) (Barjesteh et al., 2019a)

T (°C)	σmax (MPa)	σmin (MPa)	Ni	Nf
871	742.5	-532	422	487
	757	-550.4	425	474

Table 4. Results of low-cycle fatigue test on uncoated Rene^®80 specimens (ΔƐ_t (%) = 0.8, R = 0, = 2×10^-3 s^-1)

T (°C)	σmax (MPa)	σmin (MPa)	Ni	Nf
982	406	-391	155	221
	412	-369	190	253

As is seen from Tables 3 and 4, the low cycle fatigue life of the uncoated Rene^®80 decreases with increasing test temperature from 871 °C to 982 °C. The results indicated that during the test at 982 °C, specimens showed 46.6% - 54.6% lower fatigue life (N_f) compared to the specimens were tested at 871 °C. An increase in temperature led to a decreasing of the strength of γ´ phase and acceleration of an oxidation process so resulting in microstructural deterioration of the alloy (Zhang et al., 2019). The diffusion of Al into the surface as a result of oxidation at high temperature led to an increase in the area of γ´ depleted zones near the surface where initiation and nucleation of the cracks easily start from them. On the other hand, the oxide layer is more brittle than substrate, and therefore, it is an excellent site for crack formation. Rapid and frequent formation of the oxidation layer on the crack tips has an effect on the crack growth rate. This oxide layer is much more brittle than the surrounding alloy and easy to break. In the next stage, following the breakage of the oxidation layer upon the application of cyclic stress, the alloy is exposed to air and hence re-oxidized, smoothing the way for further crack growth until the final failure.

Results of low cycle fatigue test of Rene^®80 uncoated specimens at temperatures of 871 °C and 982 °C, in the form of strain vs. stress hysteresis loops are shown in Fig. 6. The hysteresis loops illustrate variations in resistance behavior of the materials against deformation under cyclic loading. Alloys may show cyclically harden, cyclically soften or a mixed behavior depending on their internal resistance to the deformation. This behavior may result from the nature and stability of the dislocation substructure of the materials (Stephens et al., 1980). As can be seen from the diagrams plotted in Fig. 6, the maximum stress obtained decreases with each cycle of strain in the low cycle fatigue test of Rene^®80 at both temperatures of 871 °C and 982 °C and at R = 0. This is known as cyclic softening. Generally in the alloy, if UTS / YS < 1.2 (where UTS is ultimate strength and YS is yield strength), the alloy will show cyclic softening behavior and if UTS /YS > 1.4, the alloy will exhibit cyclic hardening behavior (Bannantine et al., 1990). Tensile test results (Tables 1 and 2) show that the condition of UTS / YS < 1.2 is true for uncoated specimens at both temperatures of 871 °C and 982 °C. As is evident from Fig. 6a and Fig. 6b, the tension and compressive stress required to enforce the strain decrease with the increasing number of cycles.

The maximum resultant stress at the first cycle of specimen which was tested at 982 °C is lower than that of the specimen one at 871 °C. At the first cycle, initially, slip systems become active in the perpendicular direction to the cyclic stresses, and dislocations are concentrated in the γ matrix phase due to its lower strength compared to the precipitation γ´ phase. As the number of cycle increases, the deformations become homogeneous and more slip systems become active, and as a result, the density of dislocations is increased. Upon dislocation-dislocation interactions and the cutting of γ´ by dislocations, the strength of the superalloy decreases and the specimen is failed at the N_f cycle (Rahmani and Nategh, 2008b). On the other hand, by increasing the temperature, the mobility of atoms increases and diffusion-controlled mechanisms become active, dislocation mobility increases, slip becomes easier, new slip systems become available, and finally, the fracture of specimen can occur sooner (Campbell, 2012).

Figure 7 shows variations of the maximum and minimum stresses and the stress range (Δσ = σ_max – σ_min) as functions of time for uncoated specimens at two temperatures of 871 °C and 982 °C. As is evident from the Fig. 7, the rate of change of the stress reduces and the stress magnitude reaches a stable level which is named steady state condition. This condition remains stable for the rest of the fatigue life until the detection of the first fatigue crack. With the creation of the first crack at time t_i (the N_i cycle), the stress level suddenly drops and the specimen fails after a limited number of cycles at time t_f (the N_f cycle). As can be concluded from this figure, the tensile stress, decreasing at 982 °C is greater than at 871 °C, also test duration of uncoated specimen at 871 °C was about 60 min, while this time was reduced almost 46% for the same specimen at 982 °C.

Figure 8 shows the fracture surface of the uncoated specimen of the alloy Rene^®80 after low cycle fatigue failure at 871 °C and 982 °C, at the total strain range of 0.8%. Comparing Fig. 8a with Fig. 8b shows that the micro-void numbers in the superalloy substrate at higher temperature (982 °C) are more than those of lower temperature (871 °C). The presence of these micro-voids showed a ductile fracture. Also cleavage fracture (brittle fracture) can be observed in the fracture surface of both specimens. However, in comparison with specimen which was tested at 982 °C, the morphology of cleavage fracture could be seen in the larger region at 871 °C. Therefore, it could be said that the fracture mechanism was a mixed mode of ductile and brittle fracture for both temperatures, although fracture at 871 °C was more brittle. In Fig. 8c dimples at 982 °C were shown at high magnification (with increase in test temperature, profuse dimple formation occurred). Fig. 8d shows a longitudinal view of the fracture surface of the uncoated specimen, at 982 °C. The cracks were seen to be nucleated on the alloy surface under the effects of cyclic loads. As observed, the crack is nucleated on the alloy surface from the thick oxidized layer and grew in the normal direction to the loading axis. Performing tests at elevated temperature results in the surface oxidation of the uncoated alloy, and these points provide suitable sites for fatigue crack nucleation. As mentioned before, because of the diffusion of Al into the surface due to oxidation, γ´ phase will be poor in this element. Due to the depletion of the near-surface area from γ´, the alloy strength decreases, providing a basis for micro-crack nucleation. Also, more studies (Gopinath et al., 2009; Zhang et al., 2018) show that the fatigue crack nucleation has taken place beneath the surface where there are some internal defects, such as internal voids, inclusions and debris.

The typical observation of γ´ microstructure after LCF test at temperatures of 871 °C and 982 °C, is given in Fig. 9. As shown in this figure, increasing the temperature test from 871 °C to 982 °C, has no a significant effect on the γ´ morphology of the substrate, while the shape of this phase has remained almost in its cubic morphology.

3.2.3. Cyclic stress-strain behavior of coated specimens TOP

The results of the low cycle fatigue tests on the coated specimens (6 μm Pt-LTHA) at 871 °C and 982 °C are presented in Table 5. It should be mentioned that LCF test at each temperature was performed on two coated specimens.

Table 5. Results of low-cycle fatigue test on coated Rene^®80 specimens at 871 °C and 982 °C (ΔƐt (%) = 0.8, R = 0, = 2×10^-3 s^-1)

T (°C)	σmax (MPa)	σmin (MPa)	Ni	Nf
871	791	-548	386	416
	798	-554	382	410
982	400	-380	178	245
	418	-345	192	280

Experimental results show that, in coated specimens, while the total strain is equal to 0.8%, the presence of coating at 871 °C decreases the fatigue life, but in temperature of 982 °C fatigue lifetime of the coated specimens is more than that of uncoated specimens at the same total strain. At this temperature, the coated specimen tolerated 245–280 cycles to fail, indicating about 10% increase in the number of tolerated cycles, as compared to the uncoated specimen.

In Figs. 10 and 11 hysteresis loops and variations of stress with time for coated specimens at 871 °C and 982 °C are shown. As is evident from the figures, the fatigue behavior of the coated specimens resembled those of the uncoated specimen, i.e., showing a cyclic softening behavior.

It seems that, by increasing the testing temperature, the more consistency between the coating and the substrate is detected. One of the reasons supporting the increment of fatigue strength in coated specimen at 982 °C, compared to uncoated specimen, is the high ductility of this specimen. As is seen from the tensile test results (Tables 1 and 2), the coated specimen at 982 °C showed the best ductility. On the other hand, the presence of the Pt-Al coating prevents oxidation damage of the substrate alloy at elevated temperatures (Chen et al., 2015). This phenomenon plays an effective role in the limitation of the crack nucleation from the thick oxide layer which is formed on the surface of uncoated specimens. Therefore, it is expected to have a better fatigue life in coated specimen when working temperature is well above the DBTT of the coating. Figure 12 shows the surface of coated specimen after fracture at temperatures of 871 °C and 982 °C under cyclic loading (Δε_t = 0.8%).

It should be mentioned that, no difference was observed between the coated and the uncoated substrate at these temperatures, i.e., the morphology of γ´ was stable and similar to Fig. 9. A mixed mode of ductile and brittle failure was investigated at both temperatures (Fig. 12a and 12b). As seen from Fig. 12c, the cracks nucleation started from the surface of the coating at 871 °C and grew in the normal direction to the loading axis and confined to within the coating and did not penetrate the substrate. It seems that at this temperature and strain level the coating is brittle enough to create a suitable site for crack nucleation. The single phase intermediate layer of β-(Ni, Pt)Al played an important role in strength properties of the coatings and also has a marked effect on DBTT. The reasons for the brittleness of β-(Ni, Pt)Al coating (bcc structure) are the insufficient number of slip systems and grain boundary weakness (Rao, 2003). The creation of cracks in the coating causes a continuously diminishing load bearing cross-section of the coated specimen and hence the fatigue life will be declined in comparison to of uncoated specimens.

The presence of the cracks in the coating and no penetration of that to the substrate, show that, the LCF test was performed near DBTT. DBTT is related to the capability of the β-(Ni, Pt)Al phase to plastically deform under load and highly depends on chemical composition of coatings (Bose, 2007). As can be seen from Fig. 4, the weight percentage of aluminum in the coating composition was measured about 40%. This value is high and plays an effective role in the increasing of the DBTT. On the other hand, by adding the platinum in Ni-Al coating, the DBTT of the coating increases due to the improved solid solution hardening and the promotion of ductile-to-brittle phase transformations, e.g. γ→ γ’ ((Ni₃Al,Ti) (DBTT: 730–900 °C)) → β ((Ni,Pt)Al) (DBTT: 868–1060 °C)). DBTT of the PtAl₂ phase is also reported 870–1070 °C (Tamarin, 2002).

Figure 12d shows longitudinal section of the coated specimens tested at 982 °C. As can be observed from this figure, no cracks have been detected in the Pt-Al coating at this temperature. As shown in Fig. 12e, cracks initiated from internal voids and carbide interfaces at 982 °C. Result of EDS analysis at point A (Fig. 12b) has been shown in Fig. 12f. The presence of TiC as internal carbides within the dimples was confirmed by this analysis. The existence of carbide particles inside the dimples is also reported in the other research (Parlikar et al., 2017). No presence of the cracks in the Pt-Al coatings at 982 °C can be coupled to the good ductility over the DBTT. In related research (Yuan, 2013), it was mentioned that, above DBTT dislocations within the alloy are able to overcome obstacles to produce more plastic deformation. In the research performed by Rahmani and Nategh (Rahmani and Nategh, 2008b), the effect of plain-aluminide coating (CODEP-B) on low cycle fatigue properties of the Rene^®80 alloy was considered, and the results indicated that at the condition of T= 871 °C and Δεt = 0.8%, applying the coating CODEP-B on this superalloy increases LCF life. Comparing their conclusions with present results shows that Pt-Al coating on nickel-base superalloys exhibits higher DBTT than that of plain aluminides, while at T= 871 °C and Δεt = 0.8%, applying the Pt-Al coating decreases the fatigue life of Rene^®80. On the other side, by increasing the testing temperature from 871 °C to 982 °C, the LCF life of coated specimen will be improved as compared to the uncoated one tested at 982 °C. The current results indicate that the application of Pt-Al coating (6 μm Pt-LTHA) seems to have a beneficial effect on the LCF of Rene^®80 when the experimental temperature is much higher than the DBTT.

4. CONCLUSIONSTOP

REFERENCESTOP