# Oxidation Behavior and Structure Stability at 1250 °C of Chromium-Rich TaC-Containing Cast Alloys Based on Nickel and Cobalt

Patrice Berthod, Jean-Paul K Gomis, Ghouti Medjahdi

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Patrice Berthod, Jean-Paul K Gomis, Ghouti Medjahdi. Oxidation Behavior and Structure Stability at 1250 °C of Chromium-Rich TaC-Containing Cast Alloys Based on Nickel and Cobalt. Metallurgical and Materials Transactions A, 2020, 51, pp.4168-4185. 10.1007/s11661-020-05828-8. hal-04770614

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# “Oxidation Behavior and Structure Stability at 1250 °C of Chromium-Rich TaC-Containing Cast Alloys Based on Nickel and Cobalt”

Patrice Berthod, Jean-Paul K. Gomis, Ghouti Medjahdi

Metallurgical and Materials Transactions A, 2020, 51A, 4168–4185.
# Oxidation behavior and structure stability at 1250°C of chromium-rich TaC-containing cast alloys based on nickel and cobalt

Patrice Berthod<fnref n="1" />,<fnref n="2" />, Jean-Paul K. Gomis<fnref n="2" />, Ghouti Medjahdi<fnref n="1" />

**Abstract.** In this work six refractory nickel and cobalt-based cast alloys rich in chromium and designed to be strengthened by tantalum carbides were cast in controlled atmosphere. Samples were especially prepared to investigate their surface behavior in synthetic air at 1250°C for three days. Their oxidized states were first characterized by X-ray diffraction (XRD) and scanning electron microscopy (SEM) on surface prior to following metallography processing. Second, oxides and subsurfaces were examined by SEM and energy dispersion spectrometry (EDS) on cross-sections. Results show that the surface and subsurface deterioration by oxidation clearly depend on the relative quantities in Ni and Co in the chemical composition of the alloys. All alloys rather well resisted oxidation due to their chromium contents but beginnings of catastrophic oxidation were noticed for the alloys containing much more cobalt than nickel. Tantalum, present with particularly high contents, played an important role in the oxidation phenomena, notably on the adherence of the external scale, as well as on the carbides population by favoring highly stable heat-resistant TaC carbides, the fraction of which is governed by the base element.

**Keywords:** Cast alloys; Nickel and cobalt; High Ta contents; High temperature oxidation; Metallography characterization

## INTRODUCTION

<footnote n="1">¹Institut Jean Lamour, Campus ARTEM, 2 Allée André Guinier, Postal Box 50840, 54000 Nancy, France</footnote>

<footnote n="2">²Faculty of Sciences and Technologies, Campus Grignard, Rue du Jardin Botanique, Postal Box 70239, 54500 Vandoeuvre-lès-Nancy</footnote>

<footnote>patrice.berthod@univ-lorraine.fr</footnote>
Nickel and cobalt were with iron the first metals to play the role of base elements for metallic alloys considered
for the manufacturing of components destined to work at high temperature in severe conditions of mechanical
stresses and chemical aggressiveness that they may encounter in service [1]. Chromium was also often present in
their chemical compositions to form carbides for resisting creep and external oxide scales to resist oxidation by
gases and corrosion by melts [2]. The polycrystalline and equi-axed microstructures of conventionally cast nickel
and cobalt alloys allowed them to demonstrate very interesting behavior at high temperature in the mechanical field
[3]. Although advanced superalloys with very particular microstructures [4] appeared more or less recently – such as
single-crystalline alloys or Oxide Dispersion Strengthened alloys [5] – polycrystalline cast alloys reinforced by
carbides remain a good choice for many high temperature applications requiring chromium for combatting hot
corrosion. The best alloys of this category are often designed to contain MC carbides instead chromium carbides.
This is true for nickel-based alloys [7-9] as well as for cobalt-base alloys [10-12]. Although they know
morphological changes during long exposures at elevated temperatures [13, 14] the MC carbides evolve slower than
chromium carbides in the same conditions. This is the case for cobalt-based alloys containing 30wt.%Cr and TaC or
HfC, but less for NbC [15]. The HfC carbides are particularly morphologically stable at elevated temperature in
chromium-rich nickel-based [16] and cobalt-based [17] cast alloys, but this is much less true in iron-based alloys
[18]. However, taking into account both hafnium cost and availability, tantalum may be preferred to form highly
stable MC carbides.

Great quantities of TaC carbides were present as fibers in in situ composite superalloys, based on nickel [21] or
on cobalt [22], developed fifty years ago. With lower contents TaC is present as a constituent of an interdendritic
eutectic compound in the equi-axed Mar-M509 cobalt-based superalloy [1, 3, 23] in which its morphology is
Chinese script, beside the chromium carbides also present in the microstructure. More recent alloys rich in
chromium containing TaC as single carbide phase present demonstrated melting start temperatures higher than
1300°C [24] and high creep resistance at 1200°C [25]. The use of such alloys at temperatures above 1200°C may be
envisaged in case of low level of stresses applied but one can fear that catastrophic oxidation happens after only a
few hours.

The objective of this study is to test in oxidation at a temperature higher than 1200°C, for rather long duration,
several nickel-based and cobalt-based alloys rich in chromium and containing quantities of tantalum and carbon in
such proportions to obtain many TaC carbides. The compositions chosen for the alloys are defined to obtain
Co/(Ni+Co) relative amounts extended from 0 to 1 with regularly distributed values. The chromium content, which
must be high enough to favor a chromia-forming behavior but not too high to avoid significant decrease in melting
start temperature, was fixed at 25wt.%C. Since it was earlier observed that 0.4wt.%C and 6wt.%Ta were efficient to
obtain, in a cobalt-based cast alloy, script-like TaC carbides in the interdendritic spaces in quantities leading to a
good compromise between creep resistance and toughness, these values were kept here for the alloys of the present
study. So, six alloys, all containing 25Cr-0.4C-6Ta (wt.%) with a Co content increasing progressively from 0 to
about 69wt.% and a Ni content decreasing in the opposite way, were exposed in synthetic air for chosen temperature
and duration taken equal to 1250°C and 70 hours respectively.

## EXPERIMENTAL

Elaborated from pure elements (Alfa Aesar, purity > 99.9wt.%) using a high frequency induction furnace (CELES, France) in an inert gaseous environment (300 millibars of pure Argon), six ovoid ingots were obtained (weight of each: close to 40 grams). These ingots were cut to obtain first samples for the control of the as-cast microstructures and for the chemical compositions, and second flat parts for carrying out the oxidation test.

The parts destined to the alloy characterization in the as-cast state were embedded in a cold resin mixture (ESCIL, France) and ground with SiC papers from #240 to #1200. After their intermediate ultrasonic cleaning, final polishing of the embedded samples was performed with textile disk enriched with 1µm hard particles.

The parts destined to the oxidation tests were ground with #1200 SiC papers, on their main faces as well as on edges and corners to smooth them. These parts were placed in a resistive furnace continually crossed by a flow of synthetic air (80%N2-20%O2, all impurities such as water vapor being in extremely low quantity; flow: about 1.5L/h). The heating was done at a constant rate equal to 20°C/min, the isothermal stage at 1250°C during 70 hours, and the cooling rate at 5°C/min. The six samples were extracted from the cold furnace and subjected to X-ray diffraction to specify the oxides present on their surfaces.

After deposition of a thin gold layer the surfaces of the oxidized surface were examined with a Scanning Electron Microscope (JEOL, JSM-6010LA) in the Secondary Electrons mode (SEM/SE). Some spot analysis by Energy Dispersion Spectrometry (SEM/EDS) were attempted on some rare regions obviously flat and perpendicular
to the electron beam. Results were corrected from the gold contribution. X-maps were also carried out. After these
preliminary investigations, resin-embedding, cross-sectional cutting and grinding/polishing until mirror-like state
were carried out, as done in parallel for the as-cast metallographic samples.

The as-cast microstructures of the alloys were observed using the SEM in Back Scattered Electrons mode (SEM/BSE) and the chemical compositions of the alloys were controlled by EDS (except carbon). The cross-sections were observed with the SEM in BSE mode to visualize the external and internal oxides. These oxides were specified by EDS spot analyses. The same technique was used to obtain concentration profiles from the oxidation front to the bulk, perpendicularly to the oxide scale/alloy interface. The depths of the internal alloy zone affected by oxidation were measured.

RESULTS

Characteristics of the elaborated alloys

The embedded and polished samples for the as-cast state characterization were analyzed by SEM/EDS. Globally,
the wished chemical compositions were obtained, as shown in Table 1. The full frame ×250 analyses performed on
five randomly selected locations led to the presented average and standard deviation values. The standard deviations
were all lower than 1wt.%, the uncertainty of EDS analysis; consequently 1wt.% was kept for the uncertainty value.
Concerning the average values, one can see that the alloy contents in chromium were well respected. The tantalum
contents are obviously overestimated by EDS analysis: this is due to particular distribution of Ta in the alloys: the
main part of tantalum is concentrated in the TaC carbides and these ones emerge a little out of the polished surface.
This overestimation of Ta logically led to a little underestimation of nickel and cobalt. The carbon content cannot be
measured by EDS but, taking into account the obtained carbide fractions, one can guess that the wished C contents
were also successfully obtained.

A general view of the as-cast microstructures of the six alloys is given in Figure 1. One can see that all the alloys
matrixes are dendritic and that the other phases can be of two natures and that they are necessarily located in the
interdendritic spaces in the whole microstructure. The more magnified SEM/BSE micrographs displayed in Figure 2
allow distinguishing the morphologies of the interdendritic phases. The three nickel-richest alloys (named "0Co-
69Ni-Ta", "14Co-55Ni-Ta" and "27Co-42Ni-Ta") contain a double carbides population: eutectic chromium carbides (black acicular component of one of the two double-phased interdendritic compounds) and eutectic tantalum carbides (white script-like component of the other double-phased interdendritic compound). The natures of these two phase was confirmed by spot analysis on the coarsest particles found in the each sample): the Cr/C molar ratio of the black phase is close to the one of the Cr<sub>7</sub>C<sub>3</sub> carbide and the Ta/C molar ratio of the white phase is close to the one of the TaC monocarbide.

Chromium carbides are much rarer in the three cobalt-richest alloys. These alloys contain TaC as single carbide present.

**Oxidized states of the samples as characterized prior to cross-sectional preparation**

The two main faces of each of the six samples were subjected to X-ray diffraction. By alloy, the two XRD diffractograms are consistent and one of them are given in Figure 3 for the two nickel-richest alloys, in Figure 4 for the two intermediate alloys (i.e. with Ni and Co contents close to one another) and in Figure 5 for the two cobalt-richest alloys. These XRD results can be summarized by noticing that:

- The main oxide detected on the three nickel-richest alloys (Figure 3 and first diffractogram of Figure 4) is the complex oxide of chromium and tantalum, CrTaO<sub>4</sub>; by increasing the Co content in this series of three alloys chromia (Cr<sub>2</sub>O<sub>3</sub>) is more and more present; the diffraction peaks of matrix are particularly intense

- When cobalt becomes more present than nickel, the spinel oxide CoCr<sub>2</sub>O<sub>4</sub> appears and co-exists with chromia and CrTaO<sub>4</sub>; at the same time the matrix diffraction peaks become less and less high (Figure 4, Figure 5).

- For the two cobalt-richest alloys (Figure 5), cobalt oxide appears in addition to the spinel CoCr<sub>2</sub>O<sub>4</sub>, chromia and the mixed oxide CrTaO<sub>4</sub>.
<figure><img src="image_2.png" /><figcaption>**FIGURE 1.** The microstructures of the six alloys of the study, in their as-cast state (SEM/BSE, general view: ×250)</figcaption></figure>
<figure><img src="image_3.png" /><figcaption>0Co-69Ni-Ta alloy</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>14Co-55Ni-Ta alloy</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>27Co-42Ni-Ta alloy</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>42Co-27Ni-Ta alloy</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>55Co-14Ni-Ta alloy</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>69Co-0Ni-Ta alloy</figcaption></figure>

**FIGURE 2.** The microstructures of the six alloys of the study, in their as-cast state (SEM/BSE, detailed view: ×1000)
<figure><img src="image_9.png" /></figure>

<figure><img src="image_10.png" /><figcaption>FIGURE 3. X-ray diffractograms acquired on the oxidized surfaces of the six samples (for the two Ni-based alloys)</figcaption></figure>
<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /><figcaption>FIGURE 4. X-ray diffractograms acquired on the oxidized surfaces of the six samples (for the two medium alloys)</figcaption></figure>
<figure><img src="image_13.png" /></figure>

<figure><img src="image_14.png" /><figcaption>FIGURE 5. X-ray diffractograms acquired on the oxidized surfaces of the six samples (for the two Co-based alloys)</figcaption></figure>

After gold deposition the oxidized surfaces were observed using the electron microscope in secondary electrons mode. General observation allows seeing that many areas of the surfaces of the three nickel-richest alloys have lost
the oxide scale which formed during the high temperature exposure, and that the complex oxide CrTaO₄ is the most present on surface (Figure 6). Denuded alloy is less common on the surfaces of the two cobalt-richest alloys (Figure 7) while chromia and the spinel oxide are often seen (attempts of identification by EDS when possible). Some X-maps were acquired on the oxidized surfaces for completing the characterization. Two examples are showed in Figure 8 (0Co-69Ni-Ta alloy) and in Figure 9 (69Co-0Ni-Ta alloy). The first one allows clearly distinguishing the CrTaO₄ oxide covering a part of surface as well as areas of denuded alloy. On the second one a part of Cr₂O₃ scale and an area of denuded alloy are evidenced.

<figure><img src="image_15.png" /><figcaption>**FIGURE 6.** Aspect of the oxides staying on surface and of the parts of the alloys denuded by spallation (for the three Ni-richest alloys; SEM/SE, ×500)</figcaption></figure>
<figure><img src="image_16.png" /><figcaption>**FIGURE 7.** Aspect of the oxides staying on surface and of the parts of the alloys denuded by spallation (for the three Co-richest alloys; SEM/SE, ×500)</figcaption></figure>

**Oxidized states of the samples as characterized prior to cross-sectional preparation**

The surface states and the internal oxidation in subsurface as observable in cross-section are illustrated in Figure 10 by one ×1000 SEM/BSE micrograph per alloy for the four Ni-richest alloys. The external oxides formed as covering scales have obviously quitted the metallic substrate during the cooling. Oxide scale spallation was particularly severe for the Ni-richest alloys while more and more external oxides remained on surface when more and more cobalt was present in the alloy. The oxide which is the more present is the complex oxide of chromium
<figure><img src="image_17.png" /><figcaption>**FIGURE 8.** Example of a local X-ray map of the oxidized surface of the alloys (here: the cobalt-free Ni-based alloy)</figcaption></figure>

<figure><img src="image_18.png" /><figcaption>**FIGURE 9.** Example of a local X-ray map of the oxidized surface of the alloys (here: the nickel-free Co-based alloy)</figcaption></figure>
and tantalum CrTaO₄. It is mainly present as more or less interconnected subcortical islands but some CrTaO₄ films
can be observed in the parts of external scales (made of chromia mainly) remained on the alloy surface, immersed in
chromia. The external scales formed on the surfaces of the two Co-richest alloys remained more efficiently on
surface (Figure11). This allows noticing that the scales (Figure 11 top) can be rather thick (several tens of µm) and
that they may contain spinel oxide (CoCr₂O₄) and even cobalt oxide. CrTaO₄ is also present in great quantity in the
outermost part of the subsurface (Figure 11 bottom). In some locations oxidation has obviously progressed inwards
for these two cobalt-richest alloys (Figure 11 right bottom).

<figure><img src="image_19.png" /><figcaption>**FIGURE 10.** Detailed view of the oxidized surfaces of the chromia-forming alloys, with identification of the external and internal oxides by spot EDS analysis (the four Ni-richest alloys; SEM/BSE, ×1000)</figcaption></figure>
<figure><img src="image_20.png" /><figcaption>**FIGURE 11.** Detailed view of the oxidized surfaces of the chromia-forming alloys, with identification of the external and internal oxides by spot EDS analysis (the two Co-richest alloys; SEM/BSE, ×1000)</figcaption></figure>

The enlarged views ×250 presented in Figure 12 allow observing a phenomenon of disappearance of carbides (chromium carbides for the Ni-richest alloys and tantalum carbides for all alloys). The depth of carbides disappearance tends increasing from the Co-richest alloys to the Ni-richest alloys. The average depths of the carbide-free zones of alloys were measured (ten depth values per alloy) and the results are presented graphically in Figure 13. For the two Ni-richest alloys the disappearance of the chromium carbides occurred deeper than for the tantalum carbides (about 270µm against 170µm). The other alloys, which only contained initially TaC as carbide phase, they were not concerned by the chromium carbides disappearance. About the tantalum carbides which concerned all the six alloys, one can confirm that the TaC-disappearance depth is deeper for the Ni-richest alloys than for the Co-richest ones. Furthermore the TaC-depleted depth regularly decreases from the 0Co-69Ni-Ta alloy (≈180µm) to the 69Co-0Ni-Ta alloy (≈120µm).
<figure><img src="image_21.png" /><figcaption>**FIGURE 12.** View of the subsurfaces of all the alloys after oxidation (SEM/BSE, x250)</figcaption></figure>
<figure><img src="image_22.png" /><figcaption>FIGURE 13. Evolution, versus the Ni/Co ratio, of the depth of disappearance of the chromium carbides and of the one of the tantalum carbides</figcaption></figure>

Spot EDS concentration profiles were acquired from the oxide scale/alloy interface inwards the bulk. The obtained profiles are displayed in Figure 14 for the three Ni-richest alloys and in Figure 15 for the three Co-richest alloys. Concerning chromium the depth of chromium depletion is seemingly the same for all alloys (between 250 and 300µm). This depth is similar to the depth of disappearance of the chromium carbides for the two nickel-richest alloys. Such comparison cannot be done for the other alloys since they did not contain this type of carbide (or in much lower quantity, case of the 27Co-42Ni-Ta alloy). The differences concerning the chromium profiles existing between the six alloys are to be found in the outermost part of the subsurface: when the alloy contains more and more cobalt and less and less nickel the chromium content in extreme surface progressively decreases, from about from about 18 wt.% to about 12 wt.%. Logically, the average chromium content gradient becomes higher and higher, suggesting a Cr diffusion of chromium towards the oxidation front more and more difficult.. Concerning tantalum one can see in the same time a decrease in Ta content. This may be due too to more and more difficult
<figure><img src="image_23.png" /></figure>

<figure><img src="image_24.png" /></figure>

<figure><img src="image_25.png" /><figcaption>**FIGURE 14.** Concentration profiles of the main elements over a depth of 500µm for the three Ni-richest alloys after oxidation</figcaption></figure>
<figure><img src="image_26.png" /></figure>

<figure><img src="image_27.png" /></figure>

<figure><img src="image_28.png" /><figcaption>**FIGURE 15.** Concentration profiles of the main elements over a depth of 500µm for the three Co-richest alloys after oxidation</figcaption></figure>
diffusion of tantalum to be oxidized in CrTaO₄ on surface or very close to the extreme surface (as seen during the cross-sectional observations), but also to the initially higher Ta concentration in the matrix of the nickel-richest alloys resulting to the lower implication of Ta in forming carbides in their case.

Microstructures changes in the bulk due to the 70 hours stay at 1250°C

The bulk microstructures of the alloys after the exposure at high temperature are illustrated in Figure 16 (low magnification, general view) and in Figure 17 (higher magnification, detailed view). The natures, locations and volume fractions of carbides seem similar to the initial ones. In contrast their morphologies have obviously underwent coarsening and/or fragmentation. The chromium carbides of the three nickel-richest alloys have lost their acicular shape inherited from the process of eutectic solidification and they get coalesced and coarsened. The tantalum carbides of all alloys also evolved in morphology: the script-like carbides became fragmented but their coalescence and coarsening were less achieved than in the case of the chromium carbides.

DISCUSSION

Despite the low quantities of alloys (ingots of about 40g) the targeted contents in all elements were successfully obtained, thanks to the pure argon atmosphere at the pressure of several hundreds of mbars which allowed limiting oxidation and gaseous species formation. This verdict let think that not-melted parts exists in the ingots, as also suggested by the visual observation of the surfaces of all cut samples. The observed microstructures of the as-cast alloys show that solidification started by the crystallization of the FCC solid solution, with a dendritic shape. The interdenritic residual liquid solidified as eutectic {FCC matrix + carbides} compounds: {FCC matrix + Cr7C3 + TaC} ternary eutectics for the three nickel-richest alloys, and {FCC matrix + TaC} binary eutectics for the three cobalt-richest alloys. Thus, no pre-eutectic carbides precipitated prior to the matrix dendrites development for all alloys. These ones are therefore free of too coarse chromium carbide or tantalum carbides which may threaten first the mechanical resistance at high temperature and second the hot oxidation behavior by possibly disturbing locally the expected chromia-forming behavior because of a too high concentration of highly oxidable element. All carbides
<figure><img src="image_29.png" /><figcaption>0Co-69Ni-Ta alloy</figcaption></figure>

<figure><img src="image_30.png" /><figcaption>14Co-55Ni-Ta alloy</figcaption></figure>

<figure><img src="image_31.png" /><figcaption>27Co-42Ni-Ta alloy</figcaption></figure>

<figure><img src="image_32.png" /><figcaption>42Co-27Ni-Ta alloy</figcaption></figure>

<figure><img src="image_33.png" /><figcaption>55Co-14Ni-Ta alloy</figcaption></figure>

<figure><img src="image_34.png" /><figcaption>69Co-0Ni-Ta alloy</figcaption></figure>

<figure><img src="image_35.png" /><figcaption>0Co-69Ni-Ta alloy</figcaption></figure>

<figure><img src="image_36.png" /><figcaption>14Co-55Ni-Ta alloy</figcaption></figure>

<figure><img src="image_37.png" /><figcaption>27Co-42Ni-Ta alloy</figcaption></figure>

<figure><img src="image_38.png" /><figcaption>42Co-27Ni-Ta alloy</figcaption></figure>

<figure><img src="image_39.png" /><figcaption>55Co-14Ni-Ta alloy</figcaption></figure>

<figure><img src="image_40.png" /><figcaption>69Co-0Ni-Ta alloy</figcaption></figure>

<caption>FIGURE 17. The microstructures of the six alloys of the study, after 70 hours at 1250°C (SEM/BSE, detailed view: x1000)</caption>
are thus obviously of a eutectic nature and well distributed in the interdendritic spaces where they may consolidate
the dendritic structure for efficiently resisting stresses as well as the third stage of creep. However, the presence of
eutectic compounds can lead to lowered refractoriness. Fortunately, for all alloys the exposures at 1250°C did not
induce any local melting as evidencing during the bulk examination of the aged alloys, even for the three nickel-
richest alloys in which eutectic chromium carbides were present together with the tantalum carbides.

Melting did not occur in alloys during the tests at 1250°C but this temperature was obviously too high for the
morphologic stability of the carbides, especially for the chromium carbides. The natural tendency to minimize the
interface area between two solid phases (here matrix and carbide) led here to the loss of the acicular shape of the
chromium carbides which became almost round and coarse. The same tendency appeared for the tantalum carbides
but results were more limited: they get fractioned but not really coarsened. This can be maybe attributed to an
interfacial energy lower than for chromium carbides but more probably to a less easy diffusion of the big Ta atoms
in the matrix by comparison to the chromium atoms. Since one can reasonably think that this morphological
evolution of the carbides will mechanically weaken the alloys when stressed at high temperature it is clear that the
tantalum carbides are to favor at the expense of the chromium carbides. So, in the mechanical field, the cobalt-
richest alloys are advantaged by the absence of chromium carbides and by a greater proportion of tantalum carbides
by comparison to the nickel-richest alloys, in addition to the strength of the matrix which is higher when this one is
principally made of cobalt than of nickel.

Concerning oxidation it can be preliminarily noted that none of the studied alloys catastrophically oxidized despite
the very high temperature of exposure to air. All of them developed a continuous external oxide scale, even if this
was not obvious due to oxide spallation occurred during the post-isothermal stage cooling. About this problem of
oxide scale loss it can be noted that this is for the nickel-richest alloys that the spallation phenomenon was the most
intensive. This may be attributed to a possible higher participation of tantalum in the oxide formation in the
innermost part of the external scale in the case of the cobalt-free nickel-based alloy and maybe also the second Ni-
richest alloy. For the former one (0Co-69Ni-Ta alloy) no internal oxidation was noticed in the outermost part of
alloy, although all carbides, including TaC, have disappeared in the subsurface from a great depth. The
corresponding tantalum initially present as carbides has necessarily been oxidized, as external oxide since no
internal oxide has appeared. In the case of the following alloy (14Co-55Ni-Ta), which also left its carbides over a slightly lower depth than the preceding alloy, some CrTaO₄ are present over a small depth, lower than for the other alloys containing more cobalt. It can be thought that tantalum maybe diffused easier in the matrix of the two nickel-richest (flat Ta content profile) alloys than in the other alloys richer in cobalt (existing Ta content gradient), which led to localize the formation of the corresponding oxides directly on the oxidation front. This is maybe the formation of many CrTaO₄ oxides at this location that harmed the adhesion of chromia on the metallic substrate. For the other alloys, which contain more cobalt or are based on this element rather than on nickel, a slower diffusion of Ta towards the oxidation front may allow a little inwards diffusion of oxygen and therefore a CrTaO₄ formation deeper in the subsurface, as observed for the alloys with significant content in cobalt. For the later alloys, CrTaO₄ maybe thus less deleterious for the scale adhesion on alloy, with as consequence less oxide spallation at cooling for a same cooling rate.

According to the observations with the SEM in SE mode the external oxide scales were well covering the Co-rich samples and XRD runs as well as EDS measurements allowed characterizing the complex structures of the external scale formed over their surfaces. This was also true for the cross-sectional SEM observations and EDS measurements. These eternal oxide scales were constituted not only of chromia and CrTaO₄ but also of (Co,Ni)Cr₂O₄/CoCr₂O₄ spinels and (Co,Ni)O/CoO oxides, this demonstrating a behavior not wholly chromia-forming and even the imminence of catastrophic oxidation, at least locally. This is confirmed by rather irregular alloy/scale interfaces here and there for the cobalt-richest alloys. This is in agreement with the Cr contents particularly low in extreme surface by comparison with the nickel-richest alloys at the same location.

## CONCLUSION

1250°C is a very high temperature for nickel-based and cobalt-based superalloys, and particularly for the alloys studied in this work which contain eutectic compounds involving carbides. For such elevated temperature 70 hours (i.e. three days), duration very short for superalloys used at 800°C, 1000°C or even 1100°C, represent a long duration in terms of oxidation solicitation as well as of morphological stability of the carbides. Despite these severe conditions the Ni-rich versions rather well behaved in oxidation resistance. The cobalt-richest ones were much more impoverished in chromium and seemed to be ready to start to be much more damaged by oxidation. The behavior
dependence on the Co/Ni ratio maybe inversed concerning the mechanical properties at high temperature since the TaC carbides, single carbide existing in the cobalt-richest alloys and, in this way, with the greatest volume fractions for a same given carbon content, were less morphologically deteriorated than the global carbide population in the nickel-richest alloys. All these alloys seem to be exploitable at 1250°C in air for not too long periods but some of them are more interesting than the others following the oxidation/creep combinations of required properties.

## Acknowledgements

The authors wish to thank Mr. Lionel Aranda for its help for the oxidation tests.
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FIGURES CAPTIONS

**FIGURE 1.** As-cast microstructures of the six obtained alloys (SEM/BSE micrographs, large view: ×250)

**FIGURE 2.** As-cast microstructures of the six obtained alloys (SEM/BSE micrographs, enlarged view: ×1000)

**FIGURE 3.** XRD pattern of the NCT5 alloy in its as-cast state

**FIGURE 4.** Microstructures of the six alloys after 24h at 1127°C (SEM/BSE micrographs, enlarged view: ×1000)

**FIGURE 5.** Microstructures of the six alloys after 24h at 1237°C (SEM/BSE micrographs, enlarged view: ×1000)

**FIGURE 6.** Surface fraction of the chromium carbides (A) and of the tantalum carbides (B) in all alloys for all their metallurgical states (as-cast, 24h-aged at 1127C and at 1237°C)

**FIGURE 7.** Total surface fraction of the carbides in all alloys for all their metallurgical states (as-cast, 24h-aged at 1127C and at 1237°C)

**FIGURE 8.** Weight contents in chromium (A) and in tantalum (B) in the FCC Ni-based matrix of all alloys for all their metallurgical states (as-cast, 24h-aged at 1127C and at 1237°C)

**FIGURE 9.** Hardness of all alloys in all their metallurgical states (as-cast, 24h-aged at 1127C and at 1237°C); average and standard deviation from 5 indentations

**FIGURE 10.** Surface states of the six alloys after 24h at 1137°C in laboratory air (SEM/BSE micrographs, ×1000)

**FIGURE 11.** Surface states of the six alloys after 24h at 1237°C in laboratory air (SEM/BSE micrographs, ×250 and ×1000)

**FIGURE 12.** Equivalent thickness of the CrTaO₄ oxide present in subsurface close to the oxidation front, for the six alloys after 24h at 1127°C and at 1237°C in laboratory air

**FIGURE 13.** Carbide-free depth (or depth of disappearance of the BCC Cr-phase in the case of NCT5) of the six alloys after 24h at 1127°C and at 1237°C in laboratory air

**FIGURE 14.** Chromium weight content in the subsurface very close to the oxidation front, for the six alloys after 24h at 1127°C and at 1237°C in laboratory air
TABLES

<table><caption>**TABLE 1.** Chemical compositions of the obtained alloys (average value ±standard deviation, calculated from at least five SEM/EDS x250 full frame analyses).</caption><thead><tr><th>Alloys:</th><th>0Co-69Ni-Ta</th><th>14Co-55Ni-Ta</th><th>27Co-42Ni-Ta</th><th>42Co-27Ni-Ta</th><th>55Co-14Ni-Ta</th><th>69Co-0Ni-Ta</th></tr></thead><tbody><tr><th>Cr</th><td>26 ±1</td><td>26 ±1</td><td>25 ±1</td><td>26 ±1</td><td>26 ±1</td><td>25 ±1</td></tr><tr><th>Co</th><td>/</td><td>13 ±1</td><td>26 ±1</td><td>41 ±1</td><td>53 ±1</td><td>67 ±1</td></tr><tr><th>Ni</th><td>66 ±1</td><td>53 ±1</td><td>39 ±1</td><td>26 ±1</td><td>13 ±1</td><td>/</td></tr><tr><th>Ta</th><td>8 ±1</td><td>8 ±1</td><td>10 ±1</td><td>8 ±1</td><td>8 ±1</td><td>8 ±1</td></tr><tr><th>C</th><td>0.4</td><td>0.4</td><td>0.4</td><td>0.4</td><td>0.4</td><td>0.4</td></tr></tbody></table>