Tensile Behavior and Evolution of the Phases in the Al10Co25Cr8Fe15Ni36Ti6 Compositionally Complex/High Entropy Alloy

Compositionally complex alloys, or high entropy alloys, are good candidates for applications at higher temperatures in gas turbines. After their introduction, the equiatomic Al17Co17Cr17Cu17Fe17Ni17 (at.%) served as a starting material and a long optimization road finally led to the recently optimized Al10Co25Cr8Fe15Ni36Ti6 (at.%) alloy, which shows promising mechanical properties. Investigations of the as-cast state and after different heat treatments focus on the evolution of the microstructure and provide an overview of some mechanical properties. The dendritic solidification provides two phases in the dendritic cores and two different ones in the interdendritic regions. Three of the four phases remain after heat treatments. Homogenization and subsequent annealing produce a γ-γ’ based microstructure, similar to Ni-based superalloys. The γ phase is Co-Cr-Fe rich and the γ’ phase is Al-Ni-Ti rich. The understanding of the mechanical behavior of the investigated alloy is supported and enhanced by the study of the different phases and their nanohardness measurements. The observations are compared with mechanical and microstructural data from commercial Ni-based superalloys, Co-based alloys, and Co-Ni-based alloys at the desired application temperature of ~800 °C.


Introduction
High entropy alloys (single phase) and compositionally complex (more phases, based on the high entropy effect) alloys have shown promising mechanical properties since the beginning of their discovery. The current trend in the development of the new high entropy alloys (HEA) or compositionally complex alloys (CCA) is focused on their use in industrial applications as a replacement for Ni-based or Co-based alloys. The latter commonly find their applications in aircraft turbines. However, there is a high demand for increasingly efficient materials, not only for aircraft turbines, but also for the next generation of USC (ultra-supercritical) power plants. In this case, the efficiency must be increased, which means that the materials need to withstand temperatures of up to 750 • C [1]. Usually, Ni-and Co-based alloys are considered good candidates. In the temperature Tensile tests have been performed at room temperature and 800 • C to investigate the mechanical properties concerning strength and ductility. Nanohardness measurements of each phase have been carried out to support the mechanical properties.

Alloy Preparation
The Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 compositionally complex alloy (CCA) was prepared from the constituent elements of 99.99% purity. It was melted in a vacuum induction furnace and solidified in a circular ceramic mould of a 140 mm length and a 10 mm diameter. Samples were solidified directionally with a preferred [001] direction using the Bridgman method. Homogenization was carried out at 1220 • C for 20 h with a subsequent cooling in the furnace. The homogenized samples were aged at 900 • C or at 950 • C for different times, i.e., 5 and 50 h (900 • C) and 50, 100, 200, 500, and 100 h (950 • C), and then cooled down in the furnace to room temperature to prevent cracking at the grain boundaries.

Microstructural Observations
Specimens for optical microscopy (OM), scanning electron microscopy (SEM), X-ray diffraction (XRD), and nanohardness measurements were mechanically polished down in a final polishing step with a 50 nm sized OP-U colloidal silica suspension. Specimens for SEM investigation were etched with a solution of 3 g Mo-acid in 100 mL H 2 O, 100 mL HCl, and 100 mL HNO 3 .
SEM investigations were carried out in two SEMS: the first is a Zeiss Leo LEO Gemini 1530, equipped with an EDS spectrometer (Noran System Six; Thermo Fisher Scientific Inc., Waltham, MA, USA); and the second SEM is a Zeiss 1540EsB Cross Beam. All images were recorded at 5 keV and at working distances between 7.5 and 8.7 mm.
Specimens for TEM were electropolished with a solution of 83% ethanol, 10% perchloric acid, and 7% glycerin at a temperature of −7 • C and a voltage of 30 V. A TEM Philips CM30 (ThermoFischer Scientific, Waltham, MA, USA), operated at 300 kV and equipped with an energy-dispersive X-ray spectroscopy (EDS) detector (EDAX), was used in this study. The beam size for EDS measurements was 10 nm.
Phase volume fractions and precipitate sizes were determined in 2D using the training and classification of the Weka segmentation method [23] in the open source software Fiji [24] based on ImageJ [25,26]. In some cases, these measurements were accompanied by thresholds set on EDS mappings or OM, SEM, and TEM images taken along the [100] zone axis of the γ'-γ' morphology were used for the most varied view on particles and the reduction of systematic errors. An average of 1000 particles per state have been analyzed, except for the as-cast and the homogenized state, where the Weka method could not be applied due to the small size. An average of 100 particles have been observed here.

Mechanical Tests
Vickers microhardness measurements were carried out using a Polyvar met (Reichert-Jung/Leica Microsystems, Wetzlar, Germany) microscope and a MHT-10 (Anton Paar, Graz, Austria) microhardness tester under a load of~0.5 N for 15 s. The average microhardness value was calculated from the minimum ten indentations. Nanohardness measurements were performed in a picodentor HM500 (Helmut Fischer GmbH, Sindelfingen, Germany) under a load of 25 mN for 20 s.
Flat specimens for tensile tests were manufactured by electrical discharge machining (EDM) and the recast layer was removed by fine grinding. The square cross section of the samples was 1 × 1.9 mm 2 and the gauge length was 8 mm. The samples were attached form-closed to ceramic Al 2 O 3 -clampings.
For tests at elevated temperatures, a radiant heated, fire-proof furnace was used and the tensile tests were performed with a deformation rate of 0.01 mm/s. The strain data was recorded optically by a high-resolution camera and the force for deformation by a load cell. About four stress-strain value couples were logged in one second. Young's modulus was determined from the stress-strain curves by determining the slope of the elastic region.

Morphological Evolution
In a previous study [20], it has been found that the as-cast alloy solidifies dendritically, as is the case for most high entropy and compositionally complex alloys. After homogenization at 1220 • C for 20 h, the alloy does not show any dendritic structure.
In the same study, it has also been found that there are two main morphological features in the microstructure of the different states, i.e., coarse Al-Ni rich needles (in the annealed states) and a fine γ' morphology (in all states). An overview of these three phases in the four states described in [20] (as-cast, homogenized, annealed 900 • C 5 h and annealed 900 • C 50 h) is given in Figure 1. The directional solidification in the [100] direction of the samples implies that the XRD spectrum records only the (100) type planes of the γ and γ' phase and the (220) type planes of the Al-Ni rich needles, oriented 45 • towards the matrix. The latter show the strongest peak in the as-cast state, because of their high volume fraction of 10 ± 3%. The Al-Ni rich phase is absent in the homogenized state (1220 • C 20 h), and in the annealed states, its peak is visible again.

Morphological Evolution
In a previous study [20], it has been found that the as-cast alloy solidifies dendritically, as is the case for most high entropy and compositionally complex alloys. After homogenization at 1220 °C for 20 h, the alloy does not show any dendritic structure.
In the same study, it has also been found that there are two main morphological features in the microstructure of the different states, i.e., coarse Al-Ni rich needles (in the annealed states) and a fine γ' morphology (in all states). An overview of these three phases in the four states described in [20] (as-cast, homogenized, annealed 900 °C 5 h and annealed 900 °C 50 h) is given in Figure 1. The directional solidification in the [100] direction of the samples implies that the XRD spectrum records only the (100) type planes of the γ and γ' phase and the (220) type planes of the Al-Ni rich needles, oriented 45° towards the matrix. The latter show the strongest peak in the as-cast state, because of their high volume fraction of 10 ± 3%. The Al-Ni rich phase is absent in the homogenized state (1220 °C 20 h), and in the annealed states, its peak is visible again. In the following, the three phases will be presented separately. In all states of the investigated alloy, some unwanted pores and Ti-nitrides can be identified.

Al-Ni Rich Needles
These needles were not predicted by the calculation of the phase diagram in Reference [20]. They are non-existent in the homogenized state of the alloy, i.e., at 1220 °C (see [20]). They form at the lower temperatures of 900 and 950 °C. In the following, the three phases will be presented separately. In all states of the investigated alloy, some unwanted pores and Ti-nitrides can be identified.

Al-Ni Rich Needles
These needles were not predicted by the calculation of the phase diagram in Reference [20]. They are non-existent in the homogenized state of the alloy, i.e., at 1220 • C (see [20]). They form at the lower temperatures of 900 and 950 • C. The needles are oriented perpendicularly to each other. They are scarce (about 1-3 vol.%) in the 900 °C 5 h annealed alloy, but their volume fraction increases after the longer heat treatments and at a higher temperature. A comparison of the different volume fractions is shown in Figure 3. It is interesting to note that the annealing treatments at 950 °C increase the volume fraction to about 9 (100 h) or even 13% (200-1000 h). The volume fractions of needles in all states are shown in Figure 3. It can be assumed that an equilibrium is reached after 200 h at 950 °C because the volume fraction does not change after this time.
The needles' lengths remain about the same, even after the longer heat treatments at higher temperatures of 950 °C 100 h and 950 °C 1000 h, but in some cases, the needles become plate-like, especially at the grain boundaries (see Figure 2c,d).  The needles are oriented perpendicularly to each other. They are scarce (about 1-3 vol.%) in the 900 • C 5 h annealed alloy, but their volume fraction increases after the longer heat treatments and at a higher temperature. A comparison of the different volume fractions is shown in Figure 3. It is interesting to note that the annealing treatments at 950 • C increase the volume fraction to about 9 (100 h) or even 13% (200-1000 h). The volume fractions of needles in all states are shown in Figure 3. It can be assumed that an equilibrium is reached after 200 h at 950 • C because the volume fraction does not change after this time.
The needles' lengths remain about the same, even after the longer heat treatments at higher temperatures of 950 • C 100 h and 950 • C 1000 h, but in some cases, the needles become plate-like, especially at the grain boundaries (see Figure 2c,d). The needles are oriented perpendicularly to each other. They are scarce (about 1-3 vol.%) in the 900 °C 5 h annealed alloy, but their volume fraction increases after the longer heat treatments and at a higher temperature. A comparison of the different volume fractions is shown in Figure 3. It is interesting to note that the annealing treatments at 950 °C increase the volume fraction to about 9 (100 h) or even 13% (200-1000 h). The volume fractions of needles in all states are shown in Figure 3. It can be assumed that an equilibrium is reached after 200 h at 950 °C because the volume fraction does not change after this time.
The needles' lengths remain about the same, even after the longer heat treatments at higher temperatures of 950 °C 100 h and 950 °C 1000 h, but in some cases, the needles become plate-like, especially at the grain boundaries (see Figure 2c,d).

The γ-γ' Morphology in Detail
Higher resolution observations show the second most important morphological feature in the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy, i.e., the already mentioned γ-γ' morphology. As has been stated before, the microstructure of the samples aged at 900 • C and at 950 • C demonstrates that the two-phase microstructure (consisting of the γ and γ' phases) obtained in the homogenized alloy and predicted by Thermo-Calc is unstable and decomposes to form the already described additional needle-like phase. Figure 4a-d shows dark field (DF) TEM micrographs of four states of the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy, imaged using the (110) superlattice reflection. The two-phase microstructure with a γ-γ' morphology is known from Ni-based superalloys. The brightly displayed particles are γ' precipitates of an L1 2 structure embedded in a darkly displayed γ matrix of a disordered A1 structure. The corresponding diffraction pattern of the [001] zone axis is shown in the inset of every micrograph. Figure 4e,f shows two SEM micrographs of the γ-γ' morphology in the etched specimens annealed at 950 • C. These show an additional feature, i.e., small secondary γ' particles inside the γ matrix, made visible by the etching process. They will be treated in detail in a subsequent article and will not be mentioned here again. Higher resolution observations show the second most important morphological feature in the Al10Co25Cr8Fe15Ni36Ti6 alloy, i.e., the already mentioned γ-γ' morphology. As has been stated before, the microstructure of the samples aged at 900 °C and at 950 °C demonstrates that the two-phase microstructure (consisting of the γ and γ' phases) obtained in the homogenized alloy and predicted by Thermo-Calc is unstable and decomposes to form the already described additional needle-like phase. Figure 4a-d shows dark field (DF) TEM micrographs of four states of the Al10Co25Cr8Fe15Ni36Ti6 alloy, imaged using the (110) superlattice reflection. The two-phase microstructure with a γ-γ' morphology is known from Ni-based superalloys. The brightly displayed particles are γ' precipitates of an L12 structure embedded in a darkly displayed γ matrix of a disordered A1 structure. The corresponding diffraction pattern of the [001] zone axis is shown in the inset of every micrograph. Figure 4e,f shows two SEM micrographs of the γ-γ' morphology in the etched specimens annealed at 950 °C. These show an additional feature, i.e., small secondary γ' particles inside the γ matrix, made visible by the etching process. They will be treated in detail in a subsequent article and will not be mentioned here again. The primary γ' particles have different shapes and sizes in the different states. They are rather round and 40-90 nm in diameter in the as-cast state ( Figure 4a). They have undefined shapes with diameters of 6-24 nm after homogenization at 1220 • C 20 h (Figure 4b). They grow intensely after the subsequent annealing and start to adopt a cubic shape with diameters of about 200-550 nm after 900 • C 5 h (Figure 4c). They display a clear cubic shape and a diameter of 200-550 nm after 50 h annealing ( Figure 4d). The heat treatments at higher temperatures and conducted for longer times further increase the particles' sizes from 400-700 nm in the 950 • C 100 h to 400-800 nm in the 950 • C 1000 h annealed alloy.
The volume fractions of the γ' particles in all states have been summarized in Table 1 and displayed in Figure 3. They are between 30 and 55 vol.% at all states and no tendency can be determined. As-cast Homogenized 1220 • C 20 h In superalloys, the size and morphology of the γ' precipitates determine the mechanical properties. It is supposed that the same correlation can be observed in the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy.
In order to find out the growth mechanism of the γ' precipitates, especially after ageing, we compare the two-dimensional γ' particle size distribution for all states. The quantification of the particle size distribution (PSD) was made by characterizing about 100 (in the as-cast and homogenized states) or 1000 (in the annealed states) primary γ' precipitates using the following method: the area was enclosed using the Weka segmentation method and, in rare cases, additionally by hand, and the area equivalent diameter r = √ (area drawn) was obtained. The obtained values for r were then grouped into classes of 5 (as-cast), 2 (homogenized), or 50 nm (annealed). Their relative distribution in six selected states is shown in detail in Figure 5. The Gauss normal distribution (GND) is shown by a dotted line. Additionally, the PSD was calculated according to the LSW (Lifshitz-Slyozov-Wagner) theory of Ostwald ripening [27][28][29], following the equation [27]: where u = r/r avg , r is the diameter of the analyzed precipitate, and r avg is the average diameter of all precipitates. The LSW distribution is shown in a red continuous line, even for the as-cast and the homogenized sample, where no ripening is expected. The PSD in the as-cast sample (Figure 5a) spreads from about 35 to about 90 nm in diameter. After homogenization, the distribution is much narrower, i.e., from about 6 to about 24 nm in diameter (Figure 4b). For the annealed states (Figure 5c-f), the PSD is widely spread and can range from 50 to 800 nm diameter. (1) where u = r/ravg, r is the diameter of the analyzed precipitate, and ravg is the average diameter of all precipitates. The LSW distribution is shown in a red continuous line, even for the as-cast and the homogenized sample, where no ripening is expected.
The PSD in the as-cast sample (Figure 5a) spreads from about 35 to about 90 nm in diameter. After homogenization, the distribution is much narrower, i.e., from about 6 to about 24 nm in diameter (Figure 4b). For the annealed states (Figure 5c-f), the PSD is widely spread and can range from 50 to 800 nm diameter.  For all samples, the PSD fits better with the GND than with the LSW prediction. This is in accordance with research on Ostwald ripening in Ni-based alloys [30], which shows a PSD for several alloys that is broader than the one predicted by the LSW theory. A disagreement with the LSW theory has also been observed in a model Ni-14.2 Cr-5.2 Al (at.%) superalloy [31]. According to the investigations in the study from Ref. [31], the disagreement with the LSW theory can be caused by various reasons, two of which are the most plausible: (i) a transient coarsening, and (ii) the possible non-constant composition of the γ' precipitates as a function of ageing time until it reached the equilibrium values. The last argument is in agreement with the results in the present work, where the solute concentration of Al in the γ' precipitates of the annealed samples decreases with time from 14.0 at.% (after 900 • C 5 h ageing) to 11.4 at.% (after 900 • C 50 h), as determined by TEM-EDS, and shown in Table 1.
A comparison of the GND of the γ' sizes of all ten investigated states is shown in Figure 6. As has been shown in Figure 4a,b, the as-cast (purple) and the homogenized (pink) states stand out due to their small γ' sizes.
The three states after annealing at 900 • C show the most peculiar behavior. The treatment of 5 h annealing (orange dotted line) produces the largest γ' particles, but also the most wide spread of sizes. The treatment of 50 h annealing (red dotted line) shows the smallest annealing γ' precipitates, and the narrowest distribution. This is contradictory to the LSW theory, which predicts the growing of the precipitates' average radius with t 1/3 . As this result is quite unexpected, several counting methods and many micrographs have been compared to a total number of over 6000 particles of these three states and the result has been confirmed by all methods. It can be concluded that the 900 • C 50 h annealing produces a unique behavior in this Al 10  For all samples, the PSD fits better with the GND than with the LSW prediction. This is in accordance with research on Ostwald ripening in Ni-based alloys [30], which shows a PSD for several alloys that is broader than the one predicted by the LSW theory. A disagreement with the LSW theory has also been observed in a model Ni-14.2 Cr-5.2 Al (at.%) superalloy [31]. According to the investigations in the study from Ref. [31], the disagreement with the LSW theory can be caused by various reasons, two of which are the most plausible:  Table 1.
A comparison of the GND of the γ' sizes of all ten investigated states is shown in Figure 6. As has been shown in Figure 4a,b, the as-cast (purple) and the homogenized (pink) states stand out due to their small γ' sizes.
The three states after annealing at 900 °C show the most peculiar behavior. The treatment of 5 h annealing (orange dotted line) produces the largest γ' particles, but also the most wide spread of sizes. The treatment of 50 h annealing (red dotted line) shows the smallest annealing γ' precipitates, and the narrowest distribution. This is contradictory to the LSW theory, which predicts the growing of the precipitates' average radius with t 1/3 . As this result is quite unexpected, several counting methods and many micrographs have been compared to a total number of over 6000 particles of these three states and the result has been confirmed by all methods. It can be concluded that the 900 °C 50 h annealing produces a unique behavior in this Al10Co25Cr8Fe15Ni36Ti6 alloy. Unlike the annealing at 900 °C, the annealing at 950 °C produces a γ' particle growth that is in accordance with the LSW theory of particle growth (Equation (2)): where R is the average particle radius after time t, R0 is the initial particle radius, 3 is the exponent for diffusion-controlled growth, and K is the rate constant [32]. In our case, R0 corresponds to the Unlike the annealing at 900 • C, the annealing at 950 • C produces a γ' particle growth that is in accordance with the LSW theory of particle growth (Equation (2)): where R is the average particle radius after time t, R 0 is the initial particle radius, 3 is the exponent for diffusion-controlled growth, and K is the rate constant [32]. In our case, R 0 corresponds to the average particle radius after homogenization, which is around 4 nm. It can be neglected compared to the radii of the particles after annealing, and thus the equation can be simplified to Equation (3): This relationship of the average γ' particles radii versus t 1/3 and the corresponding error bars are displayed in Figure 7. It shows the linear growth of the average particle radius with the cubic root of time, with a constant of about K = 1.9.
Entropy 2018, 20, x FOR PEER REVIEW 10 of 20 average particle radius after homogenization, which is around 4 nm. It can be neglected compared to the radii of the particles after annealing, and thus the equation can be simplified to Equation (3): This relationship of the average γ' particles radii versus t 1/3 and the corresponding error bars are displayed in Figure 7. It shows the linear growth of the average particle radius with the cubic root of time, with a constant of about K = 1.9.

Mechanical Observations
The microstructure of the samples were correlated to tensile tests at room temperature and at 800 °C, in addition to Vickers micro-and nanohardness measurements.

Tensile Tests
The as-cast and the homogenized state have not been submitted to tensile tests. The tensile data from the 900 °C 5 h and the 900 °C 50 h annealed states was taken from [20]. They are shown in Table  2, together with the new tensile data of the 900 °C 50 h and the 950 °C 100 h annealed states. The engineering tensile curves of the latter two are shown in Figure 8, both at room temperature and at 800 °C.
From [20], it is known that a higher RT ultimate tensile strength σTS can be observed for the 900 °C 50 h state (1039 MPa) than for the 900 °C 5 h state (786 MPa). At a 800 °C testing temperature, their tensile strengths are almost identical (~650 MPa).
The new experiments of the 900 °C 50 h state show slightly higher RT σTS values (~1200 MPa) and lower 800 °C σTS values (~580 MPa).
The most interesting to note in Figure is the reproducibility of the tensile stress-strain curves. It is very high in the 900 °C 50 h state, but the results for the investigated 950 °C 100 h specimens are widely spread. It is almost impossible to determine an average σTS value for the investigated specimens of the 950 °C 100 h alloy. This is even worse at RT than at 800 °C.

Mechanical Observations
The microstructure of the samples were correlated to tensile tests at room temperature and at 800 • C, in addition to Vickers micro-and nanohardness measurements.

Tensile Tests
The as-cast and the homogenized state have not been submitted to tensile tests. The tensile data from the 900 • C 5 h and the 900 • C 50 h annealed states was taken from [20]. They are shown in Table 2, together with the new tensile data of the 900 • C 50 h and the 950 • C 100 h annealed states. The engineering tensile curves of the latter two are shown in Figure 8, both at room temperature and at 800 • C.
From [20], it is known that a higher RT ultimate tensile strength σ TS can be observed for the 900 • C 50 h state (1039 MPa) than for the 900 • C 5 h state (786 MPa). At a 800 • C testing temperature, their tensile strengths are almost identical (~650 MPa).
The new experiments of the 900 • C 50 h state show slightly higher RT σ TS values (~1200 MPa) and lower 800 • C σ TS values (~580 MPa).
The most interesting to note in Figure 8 is the reproducibility of the tensile stress-strain curves. It is very high in the 900 • C 50 h state, but the results for the investigated 950 • C 100 h specimens are widely spread. It is almost impossible to determine an average σ TS value for the investigated specimens of the 950 • C 100 h alloy. This is even worse at RT than at 800 • C.  A reason for this difference in fracture behavior of the 900 °C and the 950 °C annealed states can be found in the microstructure of the deformed states (see Figure ). There are three main factors that play a role in the difference in fracture behavior: 1. The deformation temperature and the implied softening of the material 2. The abundance of the Al-Ni rich needles 3. The state of the γ' precipitates A reason for this difference in fracture behavior of the 900 • C and the 950 • C annealed states can be found in the microstructure of the deformed states (see Figure 9). A reason for this difference in fracture behavior of the 900 °C and the 950 °C annealed states can be found in the microstructure of the deformed states (see Figure 9). There are three main factors that play a role in the difference in fracture behavior: 1. The deformation temperature and the implied softening of the material 2. The abundance of the Al-Ni rich needles 3. The state of the γ' precipitates There are three main factors that play a role in the difference in fracture behavior:

1.
The deformation temperature and the implied softening of the material 2.
The abundance of the Al-Ni rich needles 3.
The state of the γ' precipitates They will be compared in detail in the following, by their order of magnitude.

1.
The deformation temperature is the most important factor. The heating from RT to 800 • C implies a softening of the material and allows for a more ductile tensile behavior. The fracture micrographs of the samples deformed at RT, shown in Figure 9a,e, show little (900 • C 50 h state, Figure 9a) or no (950 • C 100 h state, Figure 9e) necking behavior and thus a quite brittle fracture. The samples deformed at 800 • C (Figure 9c,g) display an important necking and thus a ductile fracture.
This factor is no surprise and the investigated CCA behaves like many other alloys.

2.
The amount of needles is lowest in the 900 • C 5 h annealed alloy (about 1%) and highest in the 950 • C 100 h annealed alloy (about 13%). Both the longer annealing times and the higher annealing temperature increase the amount of needles, but the temperature has the stronger influence. In all five states annealed at 950 • C, the volume fraction of needles is up to ten times higher than in the 900 • C annealed states.
When observing the fracture micrographs, it was hardly possible to find any needle in the 900 • C 5 h state, and none at the fracture line (see Figure 9d). The 900 • C 50 h specimen in Figure 9b does show some needles, but the fracture seems to be uninfluenced by them because the fracture line mainly passes through the γ-γ' region. The 950 • C states in Figure 9f,h, however, show a clear fracture preference along the needles all over the specimen.
Note that the aforementioned secondary γ' particles have a preference for the regions around the needles. This is not the case in the 900 • C annealed states. This observation is helpful for determining the broken needles at the fracture line-the secondary γ' ribbons are always visible, even if the needle has quarried out.
These observations lead to the conclusion that the amount of needles is a major factor in the fracture behavior of the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy, both concerning the earlier fracture and the reproducibility of the results. Thus, in the following, the higher temperature heat treatment at 950 • C will be omitted and the next observations will be focused on the specimens annealed at the more promising heat treatments at 900 • C.
It is interesting to note that the γ' particles in the 900 • C annealed alloys (Figure 9b,d) show a heavy deformation from their original cuboid shape. They are heavily elongated in the 900 • C 50 h annealed state after deformation at RT (Figure 9b) and they adopted a starry shape in the 900 • C 5 h annealed state after deformation at 800 • C. There is hardly any deformation of the γ' particles visible in the 950 • C annealed states (Figure 9f,g). It can thus be concluded that the γ' particles absorbed the plastic deformation in the 900 • C annealed alloys, while in the 950 • C annealed alloys, the important part of the deformation energy concentrated in the interface needle-matrix.

3.
The γ-γ' morphology at room temperature has been analyzed in detail (see Section 3.1.2) and is responsible for the difference in behavior between the 900 • C 5 h and the 900 • C 50 h state. As stated above, after 900 • C 5 h annealing, the γ' precipitates have not yet established their optimum cuboidal shape. After 900 • C 50 h annealing, however, with the cuboids sides parallel to the tensile load direction and thus a higher geometrical ordering, the dislocation movement is more difficult. This effect is in accordance with the order hardening mechanism, which is often observed in Ni-based alloys [33,34].
To understand the morphological evolution of the γ' precipitates between 900 • C 5 h and 50 h ageing, the aspect ratios of the γ' precipitates were determined by using the precipitates' diameters r measured in the [100] and the [010] directions. The distribution of these aspect ratios is shown in Figure 10.
In the 900 • C 5 h annealed alloy, there is a widespread, bimodal distribution of the aspect ratios (see Figure 5c). About 37% of the analyzed precipitates have an aspect ratio below 0.75. At the same time, about 39% of them have an aspect ratio >0.9 (see Figure 10a). This result indicates that the precipitates still have no clearly defined cubic shape, as can also be seen in Figure 4c. The secondary precipitates have not been taken into account in the distribution of Figure 5c.
The 50 h annealed alloy has a narrower distribution of aspect ratios. The maximum of the distribution shown in Figure 10b is located at about 0.85-0.9. In this case, the precipitates are close to cubic.
The longer heat treatment allows the primary γ' precipitates to approach their energetically most favorable shape, i.e., cuboids, which are a compromise between the lowest surface-to-volume ratio (spheres) and the lowest elastic energy (perfect cubes). The 50 h annealed alloy has a narrower distribution of aspect ratios. The maximum of the distribution shown in Figure 10b is located at about 0.85-0.9. In this case, the precipitates are close to cubic.
The longer heat treatment allows the primary γ' precipitates to approach their energetically most favorable shape, i.e., cuboids, which are a compromise between the lowest surface-to-volume ratio (spheres) and the lowest elastic energy (perfect cubes).

Micro-and Nanohardness
The microhardness measurements were performed in the as-cast alloy in the dendritic region and in the annealed specimens randomly, including the Al-Ni rich needles. No microhardness change can be observed through the states as-cast homogenized and annealed at 900 °C. It remains 341-345 HV50/20 independently of the heat treatment. The microhardness is thus not much influenced by the presence of the Al-Ni rich needles nor by the average size of the γ' precipitates, which varies from ~14 nm (homogenized) to 312 ± 55 nm (annealed at 900 °C 5 h). The volume fraction of γ' does not contribute to the microhardness, either. The reason for the similar value of the microhardness of the samples with different states is most probably the size of the indents, which is usually about 20 μm. Thus, the measured microhardness is an average over all phases in the alloy.
The microhardness in the state annealed at 950 °C is slightly higher (about 370 compared to about 345 HV50/20 in the other states). This can probably be linked to the significantly higher amount of Al-Ni rich needles (13 vol.% compared to 1-3 vol.%).
A nanoindentation hardness map confirms that there are differences in hardness between the phases. Figure 11 shows a set of 100 nanoindents on the surface of two exemplary states, i.e., as-cast and annealed at 900 °C 50 h. The small load of 25 mN has been chosen in order to make sure that the influence of the small needles has the highest impact in the indent. Note that another set with a load of 100 mN shows less difference between the needles and the γ-γ' region and is therefore not shown here.
The as-cast alloy (Figure 11a) shows the most important local hardness changes, due to the dendritic morphology. The interdendritic regions are much harder (>700 HV0.025/20, shown in red) than the dendritic regions (~430 HV0.025/20). In the 900 °C 50 h specimen (Figure 11b), it was possible to place the indents on some Al-Ni rich needle phases (shown in red). These are harder (>500 HV0.025/20) than the γ-γ' region (~450 HV0.025/20) around them. In order to obtain better statistics, several hardness indents were placed on various Al-Ni needles (images not shown here). The 1220 °C 20 h homogenized sample (not shown here) has a very homogeneous nanohardness of 457 ± 5 HV0.025/20. The nanohardness values of the samples aged at 900 °C 5 h and 950 °C 100 h are comparable to that of the sample aged at 900 °C 50 h. All hardness values are summarized in Table  2.

Micro-and Nanohardness
The microhardness measurements were performed in the as-cast alloy in the dendritic region and in the annealed specimens randomly, including the Al-Ni rich needles. No microhardness change can be observed through the states as-cast homogenized and annealed at 900 • C. It remains 341-345 HV50/20 independently of the heat treatment. The microhardness is thus not much influenced by the presence of the Al-Ni rich needles nor by the average size of the γ' precipitates, which varies from 14 nm (homogenized) to 312 ± 55 nm (annealed at 900 • C 5 h). The volume fraction of γ' does not contribute to the microhardness, either. The reason for the similar value of the microhardness of the samples with different states is most probably the size of the indents, which is usually about 20 µm. Thus, the measured microhardness is an average over all phases in the alloy.
The microhardness in the state annealed at 950 • C is slightly higher (about 370 compared to about 345 HV50/20 in the other states). This can probably be linked to the significantly higher amount of Al-Ni rich needles (13 vol.% compared to 1-3 vol.%).
A nanoindentation hardness map confirms that there are differences in hardness between the phases. Figure 11 shows a set of 100 nanoindents on the surface of two exemplary states, i.e., as-cast and annealed at 900 • C 50 h. The small load of 25 mN has been chosen in order to make sure that the influence of the small needles has the highest impact in the indent. Note that another set with a load of 100 mN shows less difference between the needles and the γ-γ' region and is therefore not shown here.
The as-cast alloy (Figure 11a) shows the most important local hardness changes, due to the dendritic morphology. The interdendritic regions are much harder (>700 HV0.025/20, shown in red) than the dendritic regions (~430 HV0.025/20). In the 900 • C 50 h specimen (Figure 11b), it was possible to place the indents on some Al-Ni rich needle phases (shown in red). These are harder (>500 HV0.025/20) than the γ-γ' region (~450 HV0.025/20) around them. In order to obtain better statistics, several hardness indents were placed on various Al-Ni needles (images not shown here).
The 1220 • C 20 h homogenized sample (not shown here) has a very homogeneous nanohardness of 457 ± 5 HV0.025/20. The nanohardness values of the samples aged at 900 • C 5 h and 950 • C 100 h are comparable to that of the sample aged at 900 • C 50 h. All hardness values are summarized in Table 2. Unlike in the tensile tests, the preferred cuboidal shape of γ' in the 900 °C 50 h annealed alloy does not have any effect. However, the dislocation movement created by an indentation is much more arbitrary than in a directional tensile test. The orientation of the dislocation lines could thus play a role in the alloy's mechanical properties, emphasizing its anisotropic behavior. Table 2 summarizes the γ' precipitates' size and the γ' precipitates' volume fraction, the microand nanohardness, the ultimate tensile strength σTS at room temperature, and the ultimate tensile strength at 800 °C for the five mechanically investigated states.
γ' sizes and volume fractions are compared with the ultimate tensile strength at 800 °C and are summarized in Table 3, along with additional phases. Unlike in the tensile tests, the preferred cuboidal shape of γ' in the 900 • C 50 h annealed alloy does not have any effect. However, the dislocation movement created by an indentation is much more arbitrary than in a directional tensile test. The orientation of the dislocation lines could thus play a role in the alloy's mechanical properties, emphasizing its anisotropic behavior. Table 2 summarizes the γ' precipitates' size and the γ' precipitates' volume fraction, the microand nanohardness, the ultimate tensile strength σ TS at room temperature, and the ultimate tensile strength at 800 • C for the five mechanically investigated states.

Comparison with Other Alloys
The microscopic data obtained in this work allows an overview of different properties for the following alloys: Co-based Co-9Al-9W-2Ta-0.02B -Co-Ni based Co-30Ni-10Al-5Mo-2Ta and TMW-4M3-1.
The volume fraction of the γ' precipitates of the compared alloys can be classified into three groups: the highest volume fractions of about 70% are obtained in the single crystal CMSX-4, Co-based Co-9Al-9W-2Ta-0.02B and Co-Ni based Co-30Ni-10Al-5Mo-2Ta. An average volume fraction of 25-40% is obtained in the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy and Inconel 706. All other compared alloys have very low γ' volume fractions, of 5% or below.
Alloys with many additional phases, like the equiatomic Al 17 Co 17 Cr 17 Cu 17 Fe 17 Ni 17 , Inconel 706, Alloy 800, and Alloy 800H, usually have a rather low ultimate tensile strength at 800 • C. The exceptions are Inconel 617 and TMW-4M3-1. Generally speaking, a high number of phases does not seem to be advisable for an ultimate tensile strength at 800 • C.
There seems to be a correlation between the γ' volume fraction and the ultimate tensile strength at 800 • C. The latter would probably benefit from an increase of the γ' volume fraction in the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy.
Of all the eleven compared alloys, the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy shows the most similarities to the single crystal CMSX-4 alloy and the Co-9Al-9W-2Ta-0.02B alloy. The size of the γ' precipitates, the volume fraction, and the ultimate tensile strength at 800 • C lie between those of CMSX-4 and Inconel 617. Further optimization will probably continue the improvement of the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy's properties. We particularly expect in the future to improve the ultimate tensile strength at 800 • C by increasing the volume fraction of γ' precipitates by an adequate heat treatment or through the addition of traces of γ' builders.

Summary
The Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy solidifies dendritically and the dendrite cores are characterized by a two-fold morphology consisting of γ' precipitates inside a disordered fcc matrix. The dendritic structure is not present after the homogenization at 1220 • C/20 h. However, the formation of γ' precipitates could not be prevented in the homogenized sample.
After ageing at 900 • C or 950 • C, a third phase joins the γ-γ' morphology, i.e., the Al-Ni rich needles. The longer annealing times and higher annealing temperatures increase the volume fraction of the needles, from~1% after 900 • C 5 h to~13% after 950 • C 200 h. It stabilizes after 200 h at 950 • C.
The γ-γ' morphology can be observed in all states from the as-cast to the 1220 • C 20 h homogenized state and the subsequently annealed states at 900 • C and 950 • C. The size of the γ' particles is smallest after homogenization and increases drastically to over 800 nm in average diameter after the longest and hottest annealing at 950 • C 1000 h. After annealing at 950 • C, the γ' precipitates grow according to the LSW theory of Ostwald ripening. This is not the case after annealing at 900 • C. The softest annealing at 900 • C 5 h produces larger average precipitates than the annealing at 900 • C 50 h.
Tensile tests of the chosen states 900 • C 5 h, 900 • C 50 h, and 950 • C 100 h reveal that the amount of Al-Ni rich needles is the main factor responsible for an early and unpredictable fracture. The high amount in the 950 • C treated states is thus unwanted and this annealing temperature can be excluded in the future. In the 900 • C annealed samples, the longer annealing time of 50 h intensifies the cuboidal shape of the γ' precipitates, which implies better tensile results both at RT and at 800 • C. The small amount of needles does not seem to disturb the deformation mechanism.
Microhardness values are practically identical in all investigated states. At this scale, neither the γ-γ' morphology nor the Al-Ni rich needles influence the hardness. However, nanohardness maps show a difference in hardness between the softer γ-γ' region and the harder Al-Ni rich needles in the annealed alloys.
When comparing the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy to ten other alloys, i.e., Ni-based, Co-based, Co-Ni-based, and other high entropy alloys, it could be concluded that the data from the Al 10 Co 25 Cr 8 Fe 15 Ni 36 Ti 6 alloy is amongst the best in the categories γ' particle size, γ' volume fraction, and ultimate tensile strength at the desired 800 • C. The properties are approaching those of single crystal CMSX-4 and Co-based Co-9Al-9W-2Ta-0.02B at much lower costs.