Introduction

High-entropy alloys have been a focal point of research in recent years.1,2 FCC alloys based on 3d transition metals have received major research interest in search of novel materials with improved mechanical and functional properties.2,3,4,5 The most prominent example is the Co-Cr-Fe-Mn-Ni system, i.e., the ‘Cantor’ alloy CoCrFeMnNi and its derivatives.1 Recently, the impact of Cu has been addressed by adding it to the Co-Cr-Fe-Mn-Ni system, aiming to enhance mechanical properties6,7,8,9 and high-temperature wear resistance.10 These alloys are promising candidates for applications at elevated temperature. However, oxidation resistance and high-temperature microstructural stability are required11,12 to take advantage of the higher strength and ductility. It is straightforward to expect that the oxidation behavior of the respective alloys is strongly affected by microstructural features. Due to the variety of microstructures in the Co-Cr-Cu-Fe-Mn-Ni system, depending on composition and processing route, many aspects remain to be assessed.

CoCrFeMnNi is usually a single-phase FCC material, although phase separation after prolonged annealing at intermediate temperatures has been observed.13 Due to the positive mixing enthalpy when adding Cu to the other elements, the addition of Cu to the single-phase CoCrFeMnNi increases the tendency for phase separation6,7,10 and formation of Cu-rich precipitates.8 Nanoscale Cu-rich precipitates in CoCrCuFeMnNi are expected to be beneficial for mechanical properties at elevated temperatures,8 similar to Cu-containing austenitic stainless steels like Super304H.14 In some Cu-containing HEAs, particularly those containing Co, Cr, and Cu, liquid phase separation occurs, causing an inhomogeneous microstructure.15 In yet other HEAs in this system, intermetallic phases, including the σ phase, form.16,17

Oxidation of the single-phase CoCrFeMnNi alloy has been investigated in several studies.18,19,20,21 The authors generally report fast outward diffusion of Mn and the formation of non-protective, Mn-rich oxides.18,19,20,21 For Cu-containing HEAs, oxidation and microstructure formation have only sparsely been investigated, and there is no general concept of the influence of Cu on the oxidation resistance.

During air oxidation of equiatomic CoCrCuFeNi, there are reports on severe localized and inward oxidation of the Cu-rich phase.22,23,24 In the range of 800°C to 1000°C, the addition of Cu to CoCrFeNi accelerated oxidation kinetics, and the Cu-rich phase was preferentially oxidized.22 For oxidation of CoCrCuxFeNi (molar ratio x = 0; 0.5; 1.0; 1.5) at 950°C in air, an increased mass gain and deterioration of oxidation resistance with increasing Cu content were determined.25 However, substituting Mn in CoCrFeMnNi with Cu led to improved oxidation resistance between 600°C and 800°C in air.24 At the same time, a selective attack of the Cu-rich phase and the formation of ‘blisters’ was reported.24

To identify essential links between microstructure and early-stage oxidation behavior of Co-Cr-Cu-Fe-Mn-Ni high-entropy alloys, we present a systematic study investigating how composition and processing affect the microstructure formation and impact the oxidation behavior of the respective alloys. We focus on early oxidation stages to investigate the first forming oxide phases. To elucidate the influence of specific elements, the equiatomic alloy CoCrCuFeMnNi and its equiatomic sub-alloys, achieved by omitting one element each, are prepared. The microstructure is characterized along the processing steps. Early-stage oxidation in air is investigated by means of scanning electron microscopy, transmission electron microscopy, energy-dispersive x-ray spectroscopy, and x-ray diffraction. The oxidation behavior is discussed based on the experimental results.

Materials and Methods

Alloy Preparation and Processing

Seven HEAs from the Co-Cr-Cu-Fe-Mn-Ni system were prepared. Raw materials used for melting were 99.99% pure Fe and Ni (MaTeck GmbH, Jülich, Germany), 99.99% pure Cr (E. Wagener GmbH, Weissach-Flacht, Germany), 99.99% pure Cu (EVOCHEM GmbH, Offenbach a. M., Germany), 99.91% pure Mn, and 99.8% pure Co (Haines & Maassen GmbH, Bonn, Germany). Alloy compositions were chosen as equiatomic, preparing CoCrCuFeMnNi (HEA-6) and the quinary compositions achieved by omitting one alloying element, e.g., CrCuFeMnNi (Co-free HEA) by leaving out Co. The alloys were prepared using an induction levitation furnace with a cold wall crucible in Ar atmosphere of high purity (99.999 vol.%). During melting, it was ensured that high melting point metals, e.g., Cr, were fully dissolved and thorough mixing was achieved in the melt. Already at this stage it became obvious that further processing of the Ni-free and Co-free HEAs would be futile because of the largely inhomogeneous microstructure in the as-cast state. These alloys were characterized in as-cast condition without further processing.

Depending on the as-cast microstructure and expected phase formation at elevated temperature, annealing treatments were adjusted for each alloy to enable plastic deformation by rolling at room temperature. HEA-6, the Cr-free, Cu-free, and Mn-free HEAs were annealed for 72 h at 900°C and water quenched. A MgO-based coating layer was applied to prevent oxidation during the annealing treatment.26 The Fe-free HEA was annealed under Ar atmosphere for 4 h at 1000°C and water quenched. The higher annealing temperature was chosen for this alloy to prevent formation of brittle phases. All alloys beside the Co-free and the Ni-free HEAs were cold rolled to sheets of 1 mm thickness in multiple rolling steps with intermediate recrystallization annealing at 1000°C for the Fe-free HEA and 900°C for the other alloys. In the early oxidation stages investigated in the present work, the oxide thickness after oxidation did not exceed approximately 1% of the sample thickness. The final rolling step corresponded to a thickness reduction between 50% and 70%. From the rolled sheets, samples with dimensions of 10 mm × 10 mm were laser cut. The samples were ground with SiC abrasive paper successively up to 1200 grit and polished to a mirror finish using  9 µm, 3 µm, and 1 µm polycrystalline diamond suspension. After polishing, the samples were cleaned in acetone using an ultrasonic bath.

Oxidation Treatment and Characterization

Oxidation experiments were carried out for 1 h at 800°C using ambient air and a resistance muffle furnace. The short oxidation time was chosen to gain access to the early oxidation stages and the first forming oxide phases. The temperature of 800°C is in the upper region for potential applications of chromia forming alloys.11 To achieve a high heating rate, the samples were placed on a thin copper sample holder and transferred onto a pre-heated copper block in the pre-heated furnace.27 After 1 h, the samples were taken out of the furnace and left to cool in air.

The alloy microstructure was characterized using scanning electron microscopy (SEM, ZEISS EVO 40). Energy dispersive x-ray spectroscopy (EDXS) in the SEM was used to locally assess alloy and oxide compositions. Phase identification of alloy and oxide phases was carried out using x-ray diffraction (XRD, Bruker D8 Discover) with Co-Kα radiation. HEA-6, the Cr-free, Fe-free, and Mn-free HEAs after oxidation for 1 h at 800°C were mounted in epoxy to analyze cross sections of the oxide layers. The cross sections were ground and polished as described above with a final vibratory polishing step and analyzed using SEM. The oxide layer thickness after 1 h at 800°C was determined from the cross sections. The Cu-free HEA exhibited a significantly thinner oxide layer compared to the other samples. Thus, detailed characterization was carried out using transmission electron microscopy (TEM, JEOL NEOARM 200 F). Cross sections of the oxide layer on the Cu-free HEA were prepared for TEM using a dual-beam focused ion beam system (FEI Helios NanoLab 600i). Scanning transmission electron microscopy (STEM) was used to characterize the oxide layer on the Cu-free HEA. STEM-EDXS was used to map the concentration distributions in the oxide. Diffraction patterns from the oxide were recorded using nano-beam electron diffraction (NBED). Oxide phases were identified from NBED patterns and overlaid with corresponding calculated diffraction patterns using SingleCrystal 4.1.2 (CrystalMaker Software Ltd., Begbroke, UK). Thermodynamic calculations based on the CALPHAD method were carried out using FactSage 8.2 and the database ‘SGTE 2022’.28 Gibbs free energies of formation of relevant oxides at 1070 K were calculated using FactSage 8.2 and the database ‘FTOxid’.

Results and Discussion

Microstructure and Phases in the As-cast Ni-free and Co-free HEAs

The as-cast microstructure of the Ni-free (CoCrCuFeMn) and Co-free (CrCuFeMnNi) HEAs shows large-scale inhomogeneities (Fig. 1a and e) even though the melted alloys were held above their liquidus temperature in the vigorous mixing conditions of induction levitation melting. Both alloys show microstructural characteristics of liquid phase separation, indicated by the shape and arrangement of the separated Cu-poor (dark) and Cu-rich (bright) regions. The Cu-poor liquid phase apparently solidifies first, enclosing Cu-rich melt, into which dendrites of the Cu-poor phase grow during proceeding solidification.29

Fig. 1
figure 1

SEM images of the as-cast microstructure of the Ni-free HEA (a–c) and the Co-free HEA (e–g). X-ray diffractograms in the as-cast state of the Ni-free HEA (d) and the Co-free HEA (h).

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In other regions of the same samples, a microstructure without liquid phase separation, consisting of multiple phases, was found (Fig. 1b and f). For the Ni-free alloy, three phases are observed in the SEM image of this region (Fig. 1c) and identified by XRD (Fig. 1d). The phase compositions were locally measured by EDXS and are compiled in Table I. The microstructure consists of two FCC phases, i.e., a Cu-poor phase (FCC 1) and a Cu-rich phase (FCC 2). In addition, a Cr-rich σ phase is found. Similarly, two FCC phases are identified in the as-cast microstructure of the Co-free HEA—a Cu-poor phase (FCC 1) and a Cu-rich phase (FCC 2) (Fig. 1g and h). In addition, a Cr/Fe-rich BCC phase is found.

Table I Composition of phases in the as-cast microstructure of the Ni-free and Co-free HEAs as measured locally by EDXS (in at.%)
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Liquid phase separation has been reported for a number of alloys in the Co-Cr-Cu-Fe-Mn-Ni system.15,29 The phase separation is attributed to a large positive mixing enthalpy and low entropy of mixing in alloys containing Cu.15,29 To elucidate how alloy composition affects the microstructure formation (including liquid phase separation) in the present work, thermodynamic calculations were carried out. A calculated liquidus projection allows to show the primary phases or liquid phase separation. For the Ni-free HEA, a stable miscibility gap in the Cu-rich region of the quasi-ternary liquidus projection is predicted (Fig. 2a). The alloy composition (Co20Cr20Cu20Fe20Mn20) is within the miscibility gap, consistent with the experimentally observed liquid phase separation and the resulting inhomogeneous microstructure. Notably, by reducing the Cu content or adding Ni, liquid phase separation could likely be avoided. For the Ni-free HEA, liquid phase separation29 or dendritic solidification with a ‘granular dendritic phase’17 are reported in the literature. Phase identification results are consistent with the present work, indicating FCC 1, FCC 2, and σ phase formation.17,29,30

Fig. 2
figure 2

Liquidus projections displaying the primary phases or liquid phase separation for the Ni-free HEA (a) and the Co-free HEA (b). The orange dots mark the respective alloy compositions. The isopleth section along the dotted line in (b) gives the equilibrium phase diagram (black solid lines) and the metastable extension of the liquid miscibility gap (red dashed line) for the Co-free HEA (c) (Color figure online).

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For the Co-free HEA, there is a miscibility gap in the liquidus projection in the Cu-rich region (Fig. 2b); however, the alloy composition is not within this range. The red dashed line in the isopleth section in Fig. 2c shows the metastable extension of the miscibility gap in the liquid phase. When coming from the liquid state with the concentration c0, a primary phase, poor in Cu, forms. During growth of the primary phase, the melt is enriched in Cu, reaching a local minimum in the liquidus temperature. At the minimum (T1) only half of the sample is solidified. The remaining liquid is further enriched in Cu and reaches the metastable extension of the liquid miscibility gap at T2 when cooling further, leading to liquid phase separation (T3). Finally, a BCC phase and a Cu-rich FCC phase solidify, as observed in the present work. The solidification sequence is illustrated in Fig. 2c.

For the Co-free HEA, liquid phase separation has not been reported in the literature.16,17,31 Instead, dendritic solidification and microstructures consisting of FCC 1, FCC 2 and BCC phases in the as-cast state are found, consistent with the phases found in the present study.16,17,31 The largely inhomogeneous microstructures observed for the as-cast Co-free and Ni-free HEAs are expected to hinder further processing and prevent application of these alloys. Consequently, they were not processed and investigated further. Instead, we focused on alloys with single-phase or two-phase microstructure as described below.

Microstructure and Phases in As-cast and Cold-Rolled HEA-6, Cr-free, Cu-free, Fe-free, and Mn-free HEA

The as-cast microstructures of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEAs are shown in Fig. 3. Only one alloy, the Cu-free Cantor alloy, has a single-phase microstructure. The other alloys exhibit two phases, a dendritic and an interdendritic, with varying dendrite shapes and arm distances. The dendritic phase is Cu-poor and consists mostly of Co, Cr, Fe, and some Ni, appearing dark in the SEM images. The interdendritic phase is rich in Cu, Mn, and Ni, appearing bright. This is consistent with previous studies of as-cast microstructures of similar alloys.8,10,32,33,34 Small oxide particles, generally rich in Cr and/or Mn, are found in the as-cast microstructures, appearing as small black spots in the SEM images (Fig. 3). The formation of Cr/Mn-rich oxide particles in Co-Cr-Fe-Mn-Ni alloys has been reported in the literature and is presumed to originate from oxygen contamination during melting and/or from the raw materials.27,30,35

Fig. 3
figure 3

SEM images of the as-cast microstructure of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEAs.

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After annealing treatment of the as-cast alloys, they were sufficiently ductile at room temperature and were cold rolled to sheets. For the Fe-free HEA, it was determined that an annealing temperature of 1000°C is required to avoid formation of a brittle Cr-rich phase that strongly decreases ductility at room temperature.36,37 For the other alloys, annealing at 900°C is sufficient to avoid formation of brittle phases. The features in the cold-rolled microstructure are elongated along the rolling direction that is set to be horizontal in Fig. 4. The Cu-free alloy remains single-phase after annealing and cold rolling. In the Mn-free HEA and to a lesser degree in the HEA-6, nano-scale Cu-rich precipitates form inside of the dendritic phase during annealing (Fig. 4, insets). For these alloys, Cu-rich precipitate formation has been reported8,38,39 and associated with an increase in yield strength and hardness for HEA-6.8 Formation of Cu-rich precipitates was also reported for the Cr-free40 and the Fe-free HEAs.34 Interestingly, in the rolled microstructures and magnification in the SEM used here, no Cu-rich precipitates were observed for these alloys.

Fig. 4
figure 4

SEM images of the microstructure after annealing and cold rolling of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEAs. The rolling direction is horizontal in all images. The insets show Cu-rich precipitates in HEA-6 and the Mn-free HEA.

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X-ray diffraction of the cold-rolled alloys before oxidation treatment confirms two FCC phases (FCC 1 + FCC 2) for HEA-6, the Cr-free, Fe-free, and Mn-free HEAs and a single FCC phase for the Cu-free HEA (Fig. 5). Phase compositions as determined by EDXS after cold rolling are shown in Table II. The FCC 1 phase corresponds to the Cu-poor dendritic phase whereas the FCC 2 phase corresponds to the Cu-rich interdendritic phase.

Fig. 5
figure 5

X-ray diffractograms of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEA after annealing and cold rolling.

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Table II Composition of phases in the microstructure after annealing and cold rolling of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEAs as measured by EDXS (in at.%)
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Early-Stage High-Temperature Oxidation of HEA-6, Cr-free, Cu-free, Fe-free, and Mn-free HEA

In the following, the early stages of high-temperature oxidation of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEA are discussed based on experimental results. These include oxide morphology, composition, and phases determined by SEM, TEM, EDXS, and XRD. Diffractograms of the samples after oxidation for 1 h at 800°C are shown in Fig. 6. SEM images of cross sections of HEA-6, the Cr-free, Fe-free, and Mn-free HEA along with photographs of the sample surface are shown in Fig. 7 and oxide and alloy compositions obtained by EDXS in Table III. Due to the relatively thin oxide on the Cu-free HEA after 1 h at 800°C, it was investigated using TEM. The TEM results are presented in Fig. 8.

Fig. 6
figure 6

X-ray diffractograms of HEA-6, the Cr-free, Cu-free, Fe-free, and Mn-free HEA after oxidation for 1 h at 800°C. Reference reflections used are Mn3O4, Mn2O3, Fe3O4 for spinel-type oxide, CuO, and Cr2O3 for corundum-type oxide. Dashed lines are drawn to guide the eye for reflections used to unambiguously identify respective oxide phases. Black star symbols indicate reflections attributed to the FCC alloy (Color figure online).

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Fig. 7
figure 7

SEM images of cross sections of the oxide layers formed on HEA-6 (a), the Cr-free HEA (b), Fe-free HEA (c), and Mn-free HEA (d) after oxidation for 1 h at 800°C and mounting in epoxy. The positions of EDXS spot analyses are indicated in each image. The insets show photographs of the sample surface after oxidation, displaying the occurrence and extent of oxide spallation.

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Table III Composition in the cross sections at the positions indicated in Fig. 7 as measured by EDXS (in at.%)
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Fig. 8
figure 8

The Cu-free HEA after oxidation for 1 h at 800°C. STEM bright-field images of the oxide (a) and the inner oxide region (b). STEM-EDXS maps of the oxide (c). Photograph of the sample after oxidation (d). NBED pattern (black) of the outer Mn-rich oxide in zone axis [\(1\overline{1 }0\)] with overlaid calculated diffraction pattern (red) for Mn3O4 (e). High-resolution STEM annular bright-field image (f) of the Cr-rich oxide grain indicated with an arrow in (b). Corresponding FFT (black) with overlaid calculated diffraction pattern (red) for Cr2O3 in zone axis [\(\overline{4 }\overline{2 }1\)] (g). Concentration profile extracted from the EDXS maps (h); the location of the extracted concentration profile is indicated by the shaded orange area in (a) (Color figure online).

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CoCrCuFeMnNi (HEA-6)

Upon cooling after oxidation, significant spallation of the oxide on HEA-6 was observed, mostly near the sample edges (inset, Fig. 7a). The x-ray diffractogram shows reflections attributed to Mn3O4 (hausmannite, space group I4/amd) and spinel-type oxide (space group Fd-3m) (Fig. 6a). In regions where the oxide did not spall off, an oxide layer of up to 8 µm thickness is found (Fig. 7a). Many pores are visible in the oxide and along the oxide/alloy interface. The outer oxide layer is rich in Mn and Cu, likely corresponding to a Mn-Cu spinel-type oxide. The inner oxide is rich in Mn, consistent with Mn3O4. Cu is enriched in the alloy directly below the oxide (bright seam in Fig. 7a). In some regions of the sample, the Cu/Mn-rich FCC phase was not visible in the bulk below the Cu enrichment up to a depth of ~ 10 µm. This observation, together with the thick Mn-rich oxides that formed, indicates strong outward diffusion of Mn from the Cu/Mn-rich phase. Since only little Cu is incorporated in the oxide, most of it is left behind in the alloy. Some Cu can then dissolve in the Cu-poor FCC phase, creating the sub-surface zones depleted of the Cu/Mn-rich phase. Because of the limited solubility of Cu in the Cu-poor FCC phase, it additionally accumulates at the oxide/alloy interface, creating the observed Cu enrichment. Due to the selective oxidation of Mn and concomitant partial dissolution of the Cu/Mn-rich phase, the two FCC reflections are closer together than before oxidation, indicating a lattice parameter change below the surface (Fig. 6a).

CoCuFeMnNi (Cr-free HEA)

The lack of Cr in the Cr-free HEA leads to catastrophic oxidation with severe spallation of the oxide at 800°C (inset, Fig. 7b). The oxide reflections in the x-ray diffractogram (Fig. 6b) are attributed to Mn2O3 (bixbyite, space group Ia-3), spinel-type oxide, and CuO (space group C2/c). Cross sections reveal a partially detached oxide of up to 10 µm thickness with a gap between oxide and alloy (Fig. 7b). EDXS analysis indicates an oxide rich in Mn and Co, likely a Mn-Co spinel-type oxide. A small, bright seam on the oxide surface corresponds to CuO observed in the XRD measurement (Fig. 7b). However, due to the severe spallation, it is possible that only parts of the oxide layer were analyzed in the cross section. Pores are visible in the Mn/Co-rich oxide and at the oxide/alloy interface. An EDXS measurement of the alloy directly adjacent to the partially spalled oxide indicates a Cu enrichment of up to 65 at.%. This Cu enrichment is caused by the fast outward diffusion of Mn from the Mn/Cu-rich FCC phase during oxidation, because only little Cu is oxidized, similar to what was described above for HEA-6. Kirkendall pores start to form below the Cu enrichment, likely because of the fast outward diffusion of Mn.

CoCrCuMnNi (Fe-free HEA)

Similar to HEA-6, the oxide on the Fe-free HEA spalled off from parts of the sample, primarily along the sample edges (inset, Fig. 7c). The x-ray diffractogram displays reflections from Mn3O4 and spinel-type oxide (Fig. 6d). The cross-section images indicate buckling of the oxide layer and detachment from the alloy. A partially detached region of the oxide layer shows an oxide with a thickness of up to 10 µm and many pores (Fig. 7c). The EDXS analysis indicates a Mn-rich outer oxide, consistent with Mn3O4. The inner oxide is Mn/Cr-rich and likely corresponds to a Mn-Cr spinel-type oxide. Similar to the Cr-free HEA, the alloy below the oxide/alloy interface is enriched in Cu up to 65 at.%. Severe Kirkendall pore formation is observed in the Cu/Mn-rich FCC phase. This is likely due to fast outward diffusion of Mn from the Cu/Mn-rich FCC phase.

CoCrCuFeNi (Mn-free HEA)

The Mn-free HEA did not show signs of spallation after 1 h at 800°C (inset, Fig. 7d). The x-ray diffractogram shows reflections attributed to spinel-type oxide, CuO, and corundum-type oxide (space group R-3c) (Fig. 6e). Reflections from other Cu oxides, e.g., Cu2O, are not observed. The cross sections reveal inhomogeneous oxide formation for the Mn-free HEA (Fig. 7d). Depending on whether the Cu-rich phase or the Cu-poor phase was in contact with the air atmosphere, oxides of different thickness and morphology developed. In regions where the Cu-rich phase was present, an oxide of up to 5 µm thickness grew inward and outward. The outward-grown oxide contains only Cu and O and corresponds to CuO. The inward-grown oxide is rich in Fe, Cr, Co, and Ni and likely corresponds to a (Fe,Cr,Co,Ni)3O4 spinel-type oxide. The spinel-type oxide grew into the Cu-rich phase. Along the inward grown oxide, pores are observed to form at the oxide/alloy interface. In between the regions of thick oxide, a thin oxide (< 1 µm) formed where the Cu-poor phase is in contact with air. Even though this thin oxide was not investigated in detail, it likely corresponds to the corundum-type oxide detected by XRD. Due to the high Cr content and the absence of Mn and Cu in the Cu-poor phase, the formation of Cr2O3 during the early oxidation stages is plausible at these sites. This is reported in the literature for HEAs with compositions similar to the Cu-poor phase, e.g., CoCrFeNi, in some cases together with a spinel-type oxide layer.22,41 At this early stage, the Cu-rich precipitates distributed throughout the Cu-poor FCC phase do not seem to have a pronounced impact on oxidation. Previous studies have reported the formation of a thin Cr2O3 layer at the oxide/alloy interface of the inward grown oxide regions for the Mn-free HEA.23,24 In the SEM magnification used here, a similar Cr2O3 layer is not observed.

Fast outward diffusion of Cu and inward diffusion of O through the Cu-rich FCC phase contribute to the observed oxide morphology.23 CuO has a high Gibbs free energy of formation of − 120.7 kJ/mol O2 at 1070 K, whereas (Fe,Cr,Co,Ni)3O4 spinel-type oxides have a lower Gibbs free energy of formation, e.g., − 548.2 kJ/mol O2 for FeCr2O4 as calculated using FactSage. This helps to rationalize why CuO is formed on the surface, where oxygen activity is high, and the (Fe,Cr,Co,Ni)3O4 spinel-type oxide is formed below. Previous studies have noted the preferential oxidation of the Cu-rich FCC phase in CoCrCuFeNi (the Mn-free HEA) for a wide range of temperatures.22,23,24,42 Oxidation at 800°C in air led to an outer layer of CuO and intermediate and inner layers of Fe/Co/Cr-rich spinel-type oxide and Cr2O3, similar to what is observed in the present work.22,23,24 Whereas some authors also report laterally inhomogeneous oxide formation at 800°C with thicker oxides growing on the Cu-rich FCC phase,23,24 others observed a continuous oxide layer of fairly homogeneous thickness.22 In the present work, it is obvious that in the regions of the Cu-poor FCC phase, a more protective oxide can be formed during the early oxidation stages. The high Cu content in the alloy, inducing formation of a second FCC phase, prevents this oxide from covering the whole sample surface. Removing or reducing the Mn content for CoCrCuFeMnNi (HEA-6) is a straightforward way of improving the early-stage oxidation resistance, as the Mn-rich oxides containing pores no longer form, and spallation is avoided. However, preferential oxidation of the Cu-rich phase still occurs, leading to much thicker oxides in these regions.

CoCrFeMnNi (Cu-free HEA)

For the alloys discussed above, the second FCC phase led to fast outward diffusion of Mn or Cu and thus to the formation of relatively thick oxide layers prone to spallation. This does not occur for the Cu-free HEA (Cantor alloy) because of its single-phase microstructure. Even after oxidation for 1 h at 800°C, no spallation is observed (Fig. 8d). The XRD diffractogram shows reflections from Mn3O4, Mn2O3, and spinel-type oxide (Fig. 6c). Because the oxide layer after 1 h of oxidation is rather thin (< 1 µm), detailed characterization using TEM was carried out (Fig. 8). The oxide consists of an inner, ~ 200-nm-thin region of small, equiaxed grains and an outer, ~ 600-nm-thick region of larger, columnar grains (Fig. 8a and b). Pores are observed at the oxide/alloy interface. STEM-EDXS analysis reveals the multi-layered nature of the oxide (Fig. 8c and h). The outer oxide is rich in Mn with no significant content of other elements. Using NBED, the outer Mn-rich oxide is locally identified to be Mn3O4 (Fig. 8e). Even though Mn2O3 was identified on the sample using XRD, it is not found in the oxide region investigated using TEM. Due to the highly local nature of TEM, it is possible that Mn2O3 was formed in other parts of the sample surface. At the interface of the outer oxide to the inner oxide, an enrichment of Fe up to ~ 8 at.% is measured. The inner oxide region (Fig. 8b) consists of two layers. The intermediate layer, adjacent to the Mn-rich oxide, is Cr rich and contains little Mn. An oxide grain from this Cr-rich region was investigated using high-resolution STEM (Fig. 8f). A fast Fourier transform (FFT) of this high-resolution STEM image confirms a corundum-type crystal structure (Fig. 8g). This indicates that the Cr-rich layer is Cr2O3, even though no distinct corundum-type oxide reflections were detected using XRD. The inner layer, adjacent to the alloy, contains both Cr and Mn in a ratio of approximately 2:1. This is likely spinel-type MnCr2O4, corresponding to the spinel-type oxide reflections observed using XRD. Below the oxide, a distinct depletion of Mn down to essentially 0 at.% is found. The Mn depletion at the oxide/alloy interface suggests that the oxidation rate of Mn-rich oxides is limited already at this early stage by the diffusion of Mn from the bulk alloy to the oxide/alloy interface. For longer oxidation durations of 24 h to 100 h at 800°C, the diffusion of Mn through the Mn oxide was reported to be rate-limiting for oxidation of CoCrFeMnNi.19,20 It is plausible that the rate-limiting process changes from diffusion of Mn in the alloy to diffusion through the oxide for longer oxidation times as the oxide thickness increases.

Because the oxide on the Cu-free HEA contains mainly Mn and Cr, oxide stabilities of the Mn-Cr-O system are of interest. The order of oxides observed here is consistent with the Gibbs free energy of formation at 1070 K (− 507.3 kJ/mol O2 for Mn3O4, − 572.2 kJ/mol O2 for Cr2O3, − 603.6 kJ/mol O2 for MnCr2O4 as calculated using FactSage), with the innermost oxide being able to form at the lowest oxygen activity. At 800°C in air, Mn2O3, spinel-type oxides, and Cr2O3 are stable.43 These phases were all detected in the present work. In addition, Mn3O4 was observed to form at 800°C, confirmed by both x-ray and TEM diffraction, even though it is predicted to be stable only above approximately 880°C.43 In previous studies, the formation of Mn2O3 at 800°C and Mn3O4 at 900°C was observed during oxidation of the Cu-free HEA.19,20,44 However, Mn3O4 formation was observed at an even lower temperature of 800°C in 2% O2.45 Many studies focus on much longer oxidation times; therefore, it is possible that Mn3O4 formation takes place for early oxidation stages and later transforms into Mn2O3 after the transient stage approaches equilibrium.

The Cr2O3 layer forming beneath the Mn-rich oxide is not able to suppress outward diffusion of Mn.46 Fe2O3 and Cr2O3 form a solid solution;47 consequently, Fe diffusion through Cr2O3 is likely.48 Furthermore, some Fe can be incorporated in Mn3O4,45 as observed in the present case. Even though this alloy displayed the thinnest oxide among the HEAs in this study, Mn-rich oxides in combination with pores still formed. This will likely lead to a non-protective oxide that is prone to spallation at longer annealing durations,18,19 hindering high-temperature applications in air atmosphere for the alloy.

Conclusion

Microstructure formation and early-stage oxidation of equiatomic Co-Cr-Cu-Fe-Mn-Ni HEAs at 800°C were systematically investigated. The main findings are summarized as follows:

  • The Ni- and Co-free HEAs show an inhomogeneous microstructure in the as-cast state, hindering further processing. Thermodynamic calculations indicate the occurrence of liquid phase separation as the cause. For the Co-free HEA, a solidification sequence including a metastable miscibility gap in the liquid phase is proposed.

  • Alloys containing Cu exhibit a two-phase microstructure with Cu-rich and Cu-poor FCC phases. Solely the Cu-free (‘Cantor’) alloy exhibits a single-phase FCC microstructure.

  • During early oxidation in air atmosphere, particularly the combination of Cu and Mn leads to severe oxidation, specifically of the Cu/Mn-rich FCC phase. The formation of Mn-rich oxides in combination with pores and significant spallation upon cooling is observed.

  • The Mn-free HEA forms a laterally inhomogeneous oxide with a thicker oxide forming on the preferentially oxidized Cu-rich FCC phase and a thin base oxide forming on the Cu-poor phase.

  • During the early-stage oxidation, the Cu-free (‘Cantor’) alloy forms an oxide layer of the crystal structure types of Mn3O4, Mn2O3, MnCr2O4, and Cr2O3. Mn is depleted in the alloy, and Mn diffusion in the alloy is concluded to be rate limiting at this early oxidation stage.

We conclude that for high-temperature applications of Co-Cr-Cu-Fe-Mn-Ni HEAs, the formation of a Cu-rich FCC phase should be avoided and the Mn and Cu contents should be reduced to enable formation of protective oxides.