Microstructural and oxidation effects on fatigue crack initiation mechanisms in a turbine disc alloy

Effects of microstructure and oxidation on fatigue crack initiation and early propagation processes were investigated in RR1000 turbine disc alloy with different γ′ distributions and carbide distributions on the grain boundary. Fatigue tests were carried out under three-point bending and trapezoidal waveform loading (with a 90 s dwell) at 650 °C in air. The failure mode in both γ′ variants is clearly characterised by intergranular features. A number of fatigue cracks are seen to initiate at grain boundaries with bulged Co-rich oxides at the surface and/or interfaces between carbides and grain boundaries, resulting from oxidation damage assisted by applied loading. Reduced lifetime is closely linked to significant intergranular crack initiation and frequent consequent crack coalescence events, which results in enhanced fatigue crack growth (FCG) rates. The extent of intergranular features and enhanced FCG are more marked where more continuous carbides exist at the grain boundary.


Introduction
Aeroengine turbine discs operate under both elevated temperatures and cyclic loading, which can give rise to significant mechanical and environmental degradation such as oxidation fatigue, creep damage and corrosion [1]. Powder metallurgy Ni-based superalloys have been the material of choice in turbine disc applications due to their excellent resistance to such service degradation issues [2,3]. As the turbine inlet temperature of the modern gas turbine keeps rising to ensure enhanced fuel efficiency and reduced pollutant emissions, understanding the damage mechanisms during high-temperature fatigue behaviour is critical to future alloy development. Particularly oxygen-related cracking mechanisms in an oxidising environment need to be understood under these increasing temperature demands. Disc alloy performance and much of their lifetime are controlled by fatigue crack initiation processes, which are also closely associated with local microstructural features and loading condition [4,5].
In Ni-based superalloys at elevated temperatures, thermally activated processes promote multiple slip processes. Slip band cracking is less prevalent while fatigue crack initiation at the grain boundaries (intergranular cracking) is promoted more by the reduced grain boundary strength due to the diffusion of oxygen and concomitant oxidation process on the grain boundaries [6][7][8]. It is observed that a finer grain size is closely related to a greater extent of intergranular crack initiation as more grain boundaries act as diffusion paths for oxidising and embrittling species and are environmentally attacked [9]. The stress increase from oxide formation and difference in the volumes between the oxides formed and regions of grain boundaries under applied loading contributes to enhanced intergranular crack initiation, which results from the embrittlement behaviour of oxides [10,11]. It is also seen that the carbide distribution formed at the grain boundary has a significant effect on fatigue cracking behaviour under an oxidising environment [12,13]. Generally, metallic carbides (MCs) formed at a grain boundary tend to inhibit grain boundary sliding and consequent cracking behaviour at room temperature. However, MCs at the grain boundary which experience an oxidising environment can give a rise to local volume expansion and therefore enhanced misfit between the oxidised carbides and matrix [14]. This becomes more marked when a continuous formation of grain boundary carbides exists and consequently promotes intergranular crack initiation, followed by propagation along the grain boundaries, assisted by cyclic loading [14,15]. Such oxidation processes can promote enhanced crack initiation processes and multiple crack formation along the grain boundaries, which can give rise to crack coalescence events and therefore enhanced crack propagation rate. This is also linked to a greater extent of intergranular fracture features and results in reduced lifetime [16,17].
In disc alloys, fatigue fracture consists of transgranular, intergranular or a mixed mode. It is seen that fatigue testing at lower loading frequency (under dwell loading) and higher temperature in air is likely to exhibit time-dependent crack growth behaviour, where an intergranular fracture mode occurs as oxidation and/or creep effects start to interact with the fatigue failure mechanism [18,19]. This contributes to reduced FCG resistance and 1-2 order of magnitude higher FCG rate than testing in vacuum conditions [20]. An intergranular fracture mode becomes more dominant when the stress intensity factor range (DK) level (and mechanically controlled crack growth rate) is relatively low, which is associated with the available diffusion time of oxygen ahead of a crack tip per loading cycle [21,22].
In this study, fatigue tests under an oxidising environment have been conducted on RR1000 alloys having different c 0 precipitate sizes and different carbide distributions at grain boundaries produced by different heat treatment processes. Understanding the synergistic effects of microstructural features and oxidation processes on fatigue crack initiation and early growth mechanisms under extreme service conditions needs to be carefully considered in terms of life assessment. This can direct appropriate materials optimisation and heat treatment processes (i.e. controlling the size distribution of c 0 precipitate and grain size or grain boundary features together), which is closely linked to the performance and fatigue lifetime of commercial turbine disc alloys. Crack initiation and short crack behaviour at room temperature in these RR alloy variants have been discussed in our previous research, and we draw upon this for some comparisons in the current work [23]. The focus of this work is therefore on how the varying c 0 precipitate sizes and different carbide distributions at grain boundaries due to different heat treatment processes can control hightemperature fatigue initiation processes.

Materials and experimental methodology Materials
A model RR1000 alloy for turbine disc applications (a polycrystalline Ni-based superalloy) has been used for this study. The composition is presented in Table 1. The alloy was produced by a powder metallurgy route from Rolls-Royce, followed by various proprietary thermomechanical treatments and initially exhibited a bi-modal c 0 size distribution. Noncommercial heat treatments (Fig. 1a) were then conducted on the alloy to produce two different unimodal size distributions of secondary precipitates (mean sizes of 73.2 nm ± 16.3 nm and 161.4 nm ± 43 nm for the fine and coarse variants, respectively). It is also noted that the different heat treatment processes influenced the carbide distribution at grain boundaries, where a more continuous carbide distribution is seen for the coarse c 0 variant while a discrete grain boundary carbide distribution can be seen in the fine c 0 variant ( Fig. 1b and c). Figure 1 d presents electron backscatter diffraction (EBSD) mapping of the fine c 0 variant including grain distribution, grain boundary and twin boundary (the fine c 0 variant is only presented as that of the coarse c 0 variant shows a similar grain size distribution). The grain size of the two variants was similar at approximately 32 lm as both received a supersolvus solution treatment. More details about the heat treatments used and resultant microstructural features can be found in our prior work [15].
Energy dispersive X-ray (EDX) analysis of the particles on the grain boundary supports the evidence of carbide formation by representing a considerable ratio of carbon as shown in Fig. 2a although the concentration may be influenced by the surrounding matrix due to larger sampling volume of the EDX. In RR1000 alloys, the formation of the lower carbide (M 23 C 6 , Cr-Mo rich carbide) is commonly observed in the grain boundary as unstable primary carbide (MC) is expected to react with the c matrix during the heat treatment process [24]. EDX analysis ( Fig. 2b and c) conducted on such precipitates on the grain boundary after being exposed at 650°C for 10 h in the absence of load in both c 0 variants provided evidence of carbide formation and subsequent oxidation processes. It is observed that the ratio of oxygen increased with longer testing durations as expected, which is linked to more oxygen diffusion time and further progression of the oxidation reaction. Figure 3 shows how the proof stress of the fine c 0 variant varies with temperature, with a slight decrease observed between 1100 and 1000 MPa until the temperature approaches 800°C and then it drops significantly. The coarse c 0 variant follows a similar trend while a significant decrease is seen at temperatures between 500 and 600°C. The data point at 600°C requires further checking as it seems somewhat anomalous. A distinct feature of Ni-based superalloys is that yield stress does not significantly decrease with increasing temperatures until higher temperatures (e.g. 800°C). It is worth noting that both fine and coarse c 0 variants experience a significant drop in the proof stress at temperatures beyond 800°C.

Fatigue testing and characterisation
A small number of three-point bend fatigue tests were carried out on plain bend bars (dimensions * 9.3 mm 9 55 mm 9 3.8 mm). Since the temperature at the turbine disc rim is expected to be approximately 650°C, these were tested at a maximum nominal stress of 110% yield stress at 650°C, a trapezoidal waveform loading was used with 1 s rise time, 90 s dwell time at peak load and 1 s fall time to minimum load level, with a 1 s dwell at minimum load (1-90-1-1 loading) at an R-ratio of 0.1. The tests were carried out in an Instron 8501 servo-hydraulic testing machine with an ESH Ltd. high-temperature chamber attached with four high-intensity quartz IR lamps to apply heating. The temperatures of the specimens were monitored by a Eurotherm 815 thermo-controller and R-type (platinum ? 13%rhodium/platinum) thermocouple spot-welded to the specimens. Uninterrupted tests were carried out for the lifetime and fractography analysis while some tests were interrupted after certain intervals of cycles to capture different stages of fatigue crack initiation and short crack behaviour. Some fatigue tests for both c 0 variants were interrupted at around a quarter of the expected fatigue lifetime based on the uninterrupted test to capture and examine the fatigue crack initiation behaviour. Fractography analysis and post-test characterisation to understand fatigue mechanisms were carried out by optical microscope (OM), JSM 7200F Field emission scanning electron microscope (SEM) using secondary electron imaging (SEI) and backscattered electron imaging (BEI) and Oxford Aztec EDX analysis. The main focus of this work is on the mechanistic assessment through detailed fractography rather than a statistical characterisation of lifetime scatter between the variants.

Uninterrupted tests
The lifetimes at this stress level and low testing frequency of both c 0 variants are similar, 912 and 982  cycles for the fine and coarse c 0 variants, respectively. Such lifetimes are significantly lower than the lifetime of the fatigue testing under high-frequency loading at room temperature (* 150,000 cycles) [23]. This can be ascribed to the significant extent of oxygen-related damage due to the introduction of a long dwell time and resultant more diffusion time per cycle at elevated temperature. It was expected that the lifetimes of the fine c 0 variant would exhibit a higher lifetime due to the expected deleterious effect at high temperatures of the more continuous carbide formation along grain boundaries seen in the coarse c 0 variant; however, the lifetimes of both c 0 variants are similar. The overviews of the fracture surfaces of each sample can be seen in Fig. 4a (fine c 0 variant) and Fig. 5a (coarse c 0 variant). Both fracture surfaces appear quite asymmetric, but comparing these figures, the fine c 0 variant appears to have somewhat more asymmetric features, indicating the crack path around the initiation regions is more angled to the applied stress than in the coarse c 0 variant and may have occurred closer to an edge region. Figures 4 and 5 show composite overviews and close-ups of the fracture surfaces of the fine and coarse c 0 variant including crack initiation regions and propagation regions. In Fig. 4c and Fig. 5c, multiple crack initiations appear to occur (corroborated further by interrupted tests discussed later) as indicated by arrows and clear ratchet marks are also seen, all of which is indicative of multiple crack initiations at surfaces. However, it is difficult to identify where the first crack has initiated due to the large amount of intergranular crack initiation. It is seen that an intergranular fracture mode is dominant in both c 0 variants, which is ascribed to the enhanced oxidation damage on grain boundaries assisted by the introduction of a dwell time (accordingly longer   Fig. 5d. This has also been reported in our study for long FCG behaviour at elevated temperature [15].
Intergranular fracture characteristics such as a tortuous primary crack path, considerable evidence of grain boundary crack initiation and secondary crack formation are clearly discerned on the top surfaces of both c 0 variants as presented in Fig. 6. Such features are distinctive of the dwell time fatigue mechanism at elevated temperatures. The large amount of intergranular cracks and secondary cracks give rise to frequent crack coalescence events and subsequently accelerated FCG rates, which clearly contributes to the reduced lifetimes.
The specimen top surfaces were examined at higher magnification for both c 0 variants and are characterised by some distinctive intergranular fracture features as seen in Figs. 7 and 8. Thick and bulged grain boundary oxides are observed along secondary cracks and ahead of grain boundary cracking. It can be inferred that such oxides are Co-rich oxides as commonly seen in other turbine disc alloys containing a similar Co content to the RR1000 alloy [26]. It also appears that intact oxides are continuously formed along the grain boundary and remaining oxides on the secondary crack path are seen as well. It can be inferred that such intact oxides are formed before any crack initiation process and these oxides forming and then cracking contribute to crack initiation. This behaviour is more clearly discerned in the interrupted test results reported in the following section. It is notable that the cracks tend to propagate along the grain boundaries inclined normal to the applied tensile stress direction in both c 0 variants. The bulged oxides are also apt to be formed in same direction, which may indicate that the oxidation process is significantly assisted by the loading applied.
Some oxide particles remain along the cracks, which indicates grain boundary oxidation is linked to the intergranular cracking processes. In Figs. 7c and 8c, some carbide particles can be seen along the crack path as intergranular cracking has propagated through the carbides. The presence of such carbides which have formed along a grain boundary can weaken the grain boundary properties and promote crack initiation at the interface between grain boundary and carbides under an oxidising environment, and this is particularly more marked in the coarse c 0 variant with more continuous grain boundary carbide distribution. It appears that some slip bands can be also seen along the crack path as presented in Figs. 7d and 8d, and this is more marked adjacent to the final failure region, hence resulting from the high levels of plastic deformation as the crack propagated rapidly towards final failure.
EDX analysis on the bulged oxides along the grain boundaries of both c 0 variants was conducted as shown in Fig. 9. Concentration variation of Ni, Co, Al, Cr, O is presented based on the content (wt.%) across a grain boundary decorated by oxides. Enrichment in Co and O can be discerned adjacent to the grain boundary along with some content of Cr while the Ni level shows a trend of depletion at the grain boundary. This indicates that the particular bulged oxides at grain boundaries are formed by a Co-rich complex in this alloy under these testing conditions.

Interrupted tests
After interruption of the cycling at * 250 cycles (around a quarter of estimated lifetime, based on uninterrupted tests) it is seen that one crack has initiated from the edge in the fine c 0 variant, as seen in Fig. 10, while several cracks were observed at grain boundaries in the coarse c 0 variant, as shown in Fig. 11. In order to examine crack initiation behaviour more carefully, a further 200 cycles were applied to the fine c 0 variant and after this, a few more cracks were observed to have initiated at grain boundaries as well. These observations during lifetime cycling in both c 0 variants indicate that intergranular cracking is the dominant fatigue crack initiation mechanism in this alloy under this testing condition.
b Figure 4 Micrograph of the fine c 0 variant at 110% yield stress at 650°C (a) OM images showing top surfaces and fracture surface including a large ratchet mark (b) SEI image of fracture surface representing the crack initiation region and intergranular crack growth (c) SEI image of (b) at higher magnification (ratchet mark can be seen in the straight line with different contrast as marked) (d) intergranular FCG (e) extensive secondary cracks near the final failure region. It is interesting to note that the primary crack initiated from the edge in the fine c 0 variant, exhibiting a tortuous crack path and secondary crack formation as seen in Fig. 10a. Such crack features might be associated with the reduced lifetime as seen in the uninterrupted test of the fine c 0 variant. Intergranular crack initiations were also discerned along with oxides formed along grain boundaries. In the coarse c 0 variant in Fig. 11a, multiple cracks have initiated at grain boundaries along with a continuous oxide formation. It is noteworthy that a crack has initiated at the interface between the precipitate and grain boundary where a stress concentration occurs, as commonly leads to crack initiation in Ni-based superalloys. Grain boundaries are heavily decorated by oxide and carbide formation, which can promote the crack initiation and subsequent propagation. This can be attributed to the more continuous carbide formation along the grain boundaries in the coarse c 0 variant assisted by oxidation damage under 90 s dwell time at elevated temperatures.

Effects of microstructure and oxidation processes on fatigue crack initiation behaviour
It is well understood that local microstructure exerts a significant influence on fatigue crack initiation and short crack behaviour at room temperature. In terms of the expected effects of c 0 precipitate size distribution, a finer c 0 precipitate size is likely to show a dislocation shearing mechanism while coarser c 0 precipitate size tends to exhibit dislocation looping at room temperature [27,28]. Therefore, slip band crack initiation is commonly seen in turbine disc alloys, especially at lower temperatures or higher frequencies. At elevated temperatures, the crack initiation mechanism is apt to transit to intergranular crack initiation due to the significant effects of temperature and oxidation damage, although some slip band activity can be still seen especially at high strain levels. The multiple crack initiations at grain boundaries in the alloys seen in this study can be ascribed to oxidation damage, resulting in frequent crack coalescence events and significantly reduced fatigue lives.
The interface between c matrix/precipitates (i.e. grain boundary/carbide) appears to be a favourable path for oxidation and stress/strain concentration at elevated temperature as shown in Fig. 11b while it can act as effective barriers for the crack initiation and propagation at room temperature. It seems that the significant volume expansion resulting from oxidation processes (i.e. Nb carbide oxidation) is linked to the subsequent crack initiation at the interface between carbide and matrix. However, crack initiation from subsurface porosity and crystallographic faceting have been reported even at elevated temperatures at low strain levels [9]. In turbine disc alloys, the oxides that contribute to the intergranular crack initiation differ between alloys and testing conditions. In general, oxide complexes in RR1000 alloys comprise external layers of NiO/CoO and internal layers of Cr 2 O 3 /TiO 2 /Al 2 O 3 ahead of the crack tip [29]. The formation of the bulged grain boundary Ni/Co-rich oxides and Cr/ Ti/Al oxide intrusion can be seen in turbine disc alloys, and this promotes oxide cracking at elevated temperature, which results in intergranular crack initiation [26]. The impingement of slip bands at grain boundaries can be preferential paths for the diffusion of oxygen and oxide-forming elements. However, the effect of slip bands on the diffusion seems to be limited at elevated temperature compared to the grain boundaries, as the extent of the bulged Ni/Co-rich oxidation is much more marked than any evidence of slip band impingement. In this study, it also appears that bulged oxides on grain boundaries are more frequently seen than slip band formation. It has been observed that a Ni/Co-rich oxide formed along grain boundaries and ahead of the crack tip in an RR1000 alloy which has a relatively high amount of Co [29]. Such bulged Ni/Co-rich oxides are normally linked to more mismatch strain and thereby stress concentration, which can promote crack initiation process at grain boundaries and subsequent intergranular FCG behaviour.

Effects of microstructure and oxidation processes on short fatigue crack and lifetime
In terms of short crack behaviour, transgranular fracture characteristics exhibiting slip band cracking with crystallographic facets are dominant at room temperature [30][31][32]. Our previously reported room temperature testing tends to show more slip band cracking with crystallographic facet and/or porosity crack initiation [23]. On the other hand, it is known that the microstructural effects are changed at elevated temperatures due to the significant effects of environment as well as changes in the slip character. At elevated temperature, the reduced presence of facets is associated with thermally activated dislocation motion and more cross slip being promoted rather than planar slip at elevated temperatures, which can result in dominant intergranular fracture features. When the crack is significantly small, the crack driving force is relatively low, as is the crack advance rate and the fracture mode is more sensitive to the oxidation ahead of the crack tip, local microstructure, temperature and dwell time. This is pertinent to more intergranular fracture modes. On the other hand, as the crack driving force becomes greater (leading to enhanced FCG rates), the diffusion time for oxidation ahead of the crack tip becomes limited for the same increment of crack growth, which usually results in transition back to a more transgranular fracture mode.
Overall, it is seen that fatigue lifetime of both c 0 variants is significantly decreased at elevated temperature compared to lifetime at room temperature. It was expected that the fine c 0 variant would show the greater lifetime as the coarse c 0 variant exhibits a more continuous carbide distribution on grain boundaries. However, the lifetime was similar in both c 0 variants and the fine c 0 variant showed slightly more asymmetrical fracture surfaces and associated cracking. At elevated temperatures, the lifetime is closely associated with a significant number of cracks initiating at grain boundaries due to oxidation damage. This is closely linked to frequent crack coalescence events and subsequently accelerated FCG rates, offering significantly reduced lifetimes in both c 0 variants. Crack coalescence takes place even at room temperature, but is less frequent. The number of intergranular cracks is dependent on the degree of oxidation along the grain boundary as well as oxidation of the interface between c matrix/precipitate. Moreover, a continuous carbide formation is associated with weakening of the grain boundary properties and promotes crack initiation and propagation at the interface between carbides and grain boundary assisted by oxidation under cyclic loading. This can contribute to the significantly reduced lifetime of the alloys.
The number of cycles required to observe initial cracks at around a quarter of the lifetime in both c 0 variants at both room and high temperatures (although the crack in the fine c 0 variant initiated from the edge at high temperature). A similar initial crack length at room temperature is observed at 25,000 cycles and 60,000 cycles for the two fine c 0 variants and 45,000 cycles for coarse c 0 variant, which is around 2 orders of magnitude higher than the cycles to first observed crack at high temperature [23]. It should be noted that the loading condition is also different (20 Hz and 1-90-1-1 at room and high temperature, respectively). An indicative da/dN of the fine c 0 variant can be obtained by comparing the crack growth from two observations at 250 cycles (crack * 70 lm) and 450 cycles (* 150 lm) at 650°C and is estimated as * 0.4 lm/cycles. At similar crack lengths (nominal DK levels), testing at room temperature shows an average da/dN of * 0.005 lm/cycle. This indicates that short FCG rate at high temperature is around 2 orders of magnitude higher than the short FCG rate at room temperature, which can also be linked to significantly shorter lifetimes at high temperature. The accelerated short FCG rate contributes to the shortened lifetime at high temperature along with plentiful intergranular cracking (and earlier crack initiation) and subsequent crack coalescence events.

Conclusions
The effects of microstructure and oxidation on fatigue mechanisms, particularly crack initiation and lifetime, were assessed by three-point bend fatigue testing at elevated temperatures. Post-analysis of the fatigue tests was carried out by OM, SEM and EDX analysis. The summary and following conclusions are: 1. In general, the fatigue testing of both c 0 variants under 90 s dwell at elevated temperature in air is characterised by intergranular features. A number of fatigue cracks are seen to initiate at grain boundaries or/and at the interface between carbide and grain boundary with bulged Co-rich oxides at the surface that form at elevated temperatures due to the effects of dwell time (linked to diffusion of oxygen) and resultant oxidation damage. These oxides are associated with enhanced mismatch strain and therefore stress concentration, which can result in oxide cracking. 2. Fatigue lifetime of both c 0 variants is considerably reduced into the low-cycle fatigue regime assisted by the loading with a 90 s dwell time at elevated temperature compared with the lifetime at room temperature. Significantly reduced lifetime is also closely linked to significant intergranular crack initiation and frequent consequent crack coalescence events during propagation, which results in enhanced FCG rates. Moreover, the crack formation due to the oxidation of the continuous carbide distribution along grain boundaries of the coarse c 0 variant promotes accelerated FCG rates and contributes to reduced lifetimes. their help on the heat treatments and the useful discussions. For the purpose of open access, the author has applied a CC BY public copyright licence to any Author Accepted Manuscript version arising from this submission.

Authors contributions
Not applicable.

Declarations
Conflicts of interest The authors declare that they have no known competing financial interests or personal relationships.
Ethical approval Not applicable.
Open Access This article is licensed under a Creative Commons Attribution 4.0 International License, which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate if changes were made. The images or other third party material in this article are included in the article's Creative Commons licence, unless indicated otherwise in a credit line to the material. If material is not included in the article's Creative Commons licence and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder. To view a copy of this licence, visit http://creativecommons.org/licen ses/by/4.0/.