P Berthod

S Michon

S Mathieu

R Podor

C Rapin

P Steinmetz

State of the Sub-surface Microstructure of Carbides strengthened cast Superalloys after High Temperature Oxidation -Use of Thermodynamic Modelling for a better Understanding

Keywords: High Temperature Oxidation, Thermodynamic Modelling, Cast Superalloys, Carbides

Several cast strengthened superalloys, Ni base and Co base, were exposed to high temperature oxidation for long times and metallographically examined. Different phenomena occurred in the sub-surface microstructure, depending on both alloy and temperature. Thermodynamic modelling was used to know what it happened for carbon during oxidation, then to explain the observed microstructural changes. It appears that carbon atoms either quit the alloy probably after its oxidation into gases, or on the contrary go deeper into the bulk where they promote the precipitation of new carbides by solid state transformation. Thereafter, thermodynamic modelling allowed to know the new local refractoriness of the zones affected by oxidation, then to appreciate the new mechanical properties in the sub-surface.

Introduction

When exposed to high temperature oxidation most of superalloys become covered by alumina or chromia oxide scales. They are often affected by internal oxidation too, in the matrix or in grain boundaries [START_REF] Kofstad | High Temperature Corrosion[END_REF]. But other phenomena can also occur, such as microstructural modifications in the neighbourhood of the oxidation front. For carbides strengthened superalloys after high temperature oxidation, it is often seen that there is an external zone, all along the alloy surface, in which microstructure has been damaged (Fig. 1). For example, this sub-surface microstructure can encounter a disappearance of secondary carbides, a lost of the intergranular carbides with formation of holes or not, precipitation of new phases, … These modifications of the sub-surface microstructural state are not always easy to understand, notably since some elements, like carbon, cannot be analysed because of their too light atomic weight and too small weight content in the bulk. Thermodynamic modelling can be in such cases of a good help. First, when applied in an inverted sense with regard to how it is usually used, one can think that it allows to know the local content of carbon, and then what happened for this element during oxidation. Second, knowing the whole chemical composition, it leads to appreciate the consequences, on the local refractoriness, of the microstructural and chemical changes due to oxidation.

Experimental

Synthesis of the studied alloys. The alloys studied here are of two types. In order to begin this study with a simple example, a ternary Ni-Cr-C nickel alloy was first synthesised. Working on such an alloy simplifies the problem by allowing first to analyse in the best conditions the local chemical compositions (except carbon) without any interaction between two elements, and second to perform thermodynamic calculations using a well known database. It was prepared in small quantity (ingots of nearly 50g) by high frequency induction melting under argon gas in a copper cold crucible in which it also solidified. Its microstructure is constituted of a nickel austenitic matrix (including a part of the Cr and C atoms in solid solution) and of chromium carbides at grain boundaries. Second, two cobalt base superalloys were obtained using another induction furnace (capacity of several kg) under argon gas too. The molten alloys were poured and solidified in a sand mould. Thereafter these two cobalt base superalloys underwent a homogenisation treatment at 1200°C for 2 hours, followed by a precipitation stage at 1000°C for 10 hours, in order to get a precipitation of thin carbides in matrix. These Co base superalloys contain mainly cobalt, nickel and chromium but also tantalum and carbon, the latter elements promoting the formation of refractory Cr 23 C 6 and/or TaC carbides which are known to be of a great importance for mechanical properties at high temperature [START_REF] Bradley | Superalloys : A technical guide[END_REF][START_REF] Sims | The superalloys[END_REF]. to become electrically conductive. Then a sufficiently thick electrolytic nickel coating was deposited in order to protect the external chromia layer before cutting. Thereafter samples were cut, embedded in epoxy resin and polished up to 1200 grid paper. A Scanning Electron Microscope (SEM) Philips XL30 was used in Back Scattered Electrons (BSE) mode for the surface and bulk examinations. The constitutive phases can be clearly evidenced since their average atomic numbers are sensibly different. In addition, some of them were identified by Energy Dispersion Spectrometry (EDS), or by Wavelength Dispersion Spectrometry (WDS) using a Cameca SX100 microprobe. EDS enabled also to analyse the mean chemical composition of a selected area over the local microstructure while microprobe was also used to get concentration profiles through the external zone damaged by oxidation.

Alloys

Image analysis and thermodynamic calculations. Surface fractions of the different phases were estimated on SEM photographs (BSE mode), using the AnalySIS software of Soft Imaging System GMBH. Making the hypothesis that surface fraction and volume fraction have similar values, volume fractions were assessed. Using the equation (Eq. 1) with room temperature densities of the different phases, it was possible to estimate the mass fraction of each phase. The values used for the specific weights (in g.cm -3 ) were 8.90 for the matrix, 6.94 for the Cr 7 C 3 carbide, 6.95 for the Cr 23 C 6 carbide and 14.50 for the TaC carbide [START_REF][END_REF].

∑ ϕ ρ ϕ ρ = ϕ ϕ ϕ i i vol j vol j mass )) ( f . ( ) ( f . ) ( f i j (1)
Calculations on the simple nickel alloy were carried out using Windows version Thermo-Cal software with the SSOL database (SGTE). For the cobalt alloys, N Thermo-Calc software was used in association with a new database. This last one is the SSOL database completed with binary and ternary systems involving the Ta element. It had been already used to study analogous alloys in other studies [5,6,7].

Results obtained on the simple nickel alloy (Alloy1)

Surface microstructure after oxidation at 1000°C and 1100°C for 100 hours. At 1000°C, the initial sub-surface microstructure of Alloy1 is transformed into two separated zones (Fig. 2). The outer one, thereafter called "carbideless zone", has completely lost its carbides. The inner zone still contains some carbides but with a rounder shape. One also can see a population of new phases with an acicular shape, different sizes and precise orientations. Spot microanalysis showed that both acicular particles and compact particles are Cr 7 C 3 carbides. WDS chromium profiles, were compared to the different microstructural zones. The carbideless zone corresponds to a high Cr gradient and a chromium range of [21%wt , 26%wt] while the zone containing acicular carbides exhibits a lower Cr gradient and higher Cr contents. SEM photographs were taken on several areas (of about 80µm x 120µm) in the acicular precipitation zone, on which image analysis were performed. EDS analysis of the chemical composition was realised on the same areas. The six results lead to an average of 4.6%mass of Cr 7 C 3 carbides (mass fraction) in the microstructure and to a Cr content of 29.8%wt Cr. At 1100°C, the damaged zone of Alloy1 seems to be only composed of the carbideless zone, which is three times deeper than the one resulting from oxidation at 1000°C. From the bulk, the WDS Cr content profile shows that chromium content decreases from 31%wt Cr to about 25%wt Cr. Thermodynamic calculations. An isothermal ternary Ni-Cr-C diagram was computed for T=1000°C (Fig. 3). Considering the nickel-rich corner and the Cr contents measured with WDS On the isothermal diagram obtained for 1100°C, the absence of any other phases than matrix indicates that the carbon content must be lower than 0.054%wt C at the extreme surface (where %wtCr=25) and lower than 0.025%wt C at the junction of the carbideless zone and the bulk (where %wtCr=31). In this way, this zone of the alloy has lost a great part of its initial carbon content.

Everywhere else in the alloy, the carbide surface fraction seems to be the same as in the bulk, then there is no region that became enriched with a part of carbon lost by the damaged zone.

Results obtained on the two cobalt base superalloys (Alloy2 and Alloy3) Surface microstructure after oxidation at 1000°C and 1200°C for 100 hours. Initial bulk microstructure is composed first of an austenitic cobalt matrix with other elements in solid solution.

In the interdendritic spaces of Alloy2 there are tantalum carbides belonging to an eutectic with the matrix, whereas Alloy3 contains both eutectic TaC carbides and eutectic Cr 7 C 3 in grain boundaries. These carbides strengthen these superalloys which are very resistant mechanically at high temperature. Apart the external (Cr 2 O 3 ) and internal (Ta 2 O 5 or CrTaO 4 ) oxides, new acicular oriented phases appeared in the alloys not far from the oxidation front. They are visible in a zone separating a thin carbideless zone and the not damaged zone (Fig. 4).

examples of zones on which surface fractions were measured

Alloy2

Alloy3 zone, which is of about 15µm for Alloy2 and of 10µm for Alloy3. EDS profiles were performed through the carbideless zone. They showed for Alloy2 that the chromium and tantalum contents decrease respectively from 20%wt Cr to 16%wt Cr and from 1.0%wt Ta to 0.0 %wt Ta between the frontier of the zone which still contains carbides and the extreme surface. For Alloy3, these contents decrease from 21%wt Cr to 18%wt Cr and from 1.6%wt Ta to 1%wt Ta. Several high magnification SEM photographs were taken on delimited areas of about 500µm 2 in the acicular precipitation zone (Fig. 4 right) to be thereafter subjects to image analysis measurements.

On the same areas global chemical compositions were analysed by EDS. For each alloy three results were obtained and they showed similar values. On these areas Alloy2 contains 3.3%surf Alloy3 was also exposed to oxidation at 1200°C for 100 hours. Apart an external chromia layer three times thicker and an internal oxidation of Ta more developed than for 1000°C, it exhibits a deeper carbideless zone (Fig. 5). Like for Alloy1 at 1100°C, no precipitation of new phases is visible. Here too EDS spot analysis were performed across the carbideless zone. From the frontier of the not damaged zone towards the extreme surface, Cr and Ta contents decrease from 22%wt Cr and 1.5%wt Ta to 17%wt Cr and 0.3%wt Ta, the nickel content staying at about 10%wt Ni.

Fig. 5: Surface and sub-surface microstructural state of Alloy2 oxidised at 1200°C for 100 hours Thermodynamic calculations. For Alloy2 oxidised at 1000°C and for the EDS results obtained, Thermo-Calc N indicates that the carbon weight content must be lower than 0.076%wt C at the extreme surface and lower than 0.050%wt C at the other side of the carbideless zone. These values are about ten times smaller than the initial C content of Alloy2. For Alloy3 at the same locations, calculations show that carbon content is smaller than 0.061%wt C and 0.019%wt C respectively, as is to say between ten and twenty times smaller than the initial C content of this alloy. Calculations were also performed on the compositions of the zones containing acicular carbides. For Alloy2, calculations converged to the experimental mass fractions of phases resulting from image analysis for 0.27%wt C. The final calculation results gave 2.6%mass of Cr 23 C 6 and 1.6%mass of TaC. For Alloy3, calculations allowed to come close to the real mass fractions for 0.55%wt C. For Alloy2 oxidised at 1200°C, from the extreme surface towards the beginning of the bulk the maximal possible carbon content decreases from 0.23%wt C down to 0.08%wt C. Like Alloy1 at 1100°C, no zone enriched with carbon can be seen here all over the bulk. In the carbideless zone and in the acicular carbides zone, the new local microstructure is due to an evolution of the chemical composition. It can induce changes for the local refractoriness of each zone in the alloy, and there is a melt hazard if working temperature suddenly raises quickly, as it can be encountered for hot industrial pieces. Using ThermoCalc on the local new chemical compositions allows to know what are locally the new solidus temperature. Table 2 displays these calculated values with two cases for the carbideless zones : carbon being equal to its lowest possible value (0%) or to its highest one (C max ). In fact it appears generally that these new microstructural areas are more refractory than the initial alloy, except for the acicular carbide zone (the solidus of which is sensibly equal to the bulk one).

Alloys

Interpretation and discussion

The existence of the carbideless zone is possible only if carbon quantity is smaller than very low values. There, carbon content can be considered as negligible compared to carbon contents of initial alloys and of neighbour zones containing old or new carbides. This zone has almost lost all its carbon atoms whereas the zone in which acicular carbides appeared exhibits a higher carbon content that the initial one. Since bulk presents the initial carbon content, the additional carbon necessary comes from the zone which has lost its carbides. It is possible to simplify the carbon content profile through the damaged zone (Fig. 6). Taking the example of the Alloy1 oxidised at 1000°C, it can be considered that the carbon content is zero across the carbideless zone. It increases quickly and stays at 0.435% wt C across the acicular precipitation zone, and falls quickly to the bulk content 0.20% and thereafter stays at this value. The mass of carbon in the damaged zone is equal to (S.ρ).(e 1 .%wtC 1 + e 2 .%wtC 2 ). S is the considered surface, ρ is the alloy density (about 8.9g.cm -3 ). e i and %wtC i are the depths and carbon contents of the two parts of the damaged zone. For 1cm 2 of external surface, we find about 0.30mg of C in the damaged zone, which is sensibly the initial carbon mass contained in the same volume (0.27mg). Then, almost all the carbon lost by the carbideless zone came into the matrix more deeper in the alloy. This carbon left the grain boundary chromium carbides, the decomposition of which was promoted by the diffusion of Cr towards the oxidation front. Thereafter, they diffuse backwards and entered into the matrix and met chromium atoms present in solid solution. There carbon formed with chromium new carbides which precipitated in matrix with an acicular shape. In a same grain, these new carbides exhibit a same orientation. They may coexist with the initial grain boundary carbides which have not yet disappeared and which have a more compact shape. At 1100°C, the mechanisms would not be the same, since no carbon enrichment can be noted elsewhere in the alloy. Carbon lost by decomposition of the initial grain boundary carbides disappears, without any doubt being oxidised into CO or CO 2 gases. For 1cm 2 of Alloy1 external surface and a carbideless zone of about 100µm in which carbon content is negligible, the same type of calculation as before leads to a carbon mass lost of 180mg that became gas. Unfortunately, such calculations are not really possible for the two cobalt alloys. Indeed, zones with acicular precipitated carbides often are partially mixed in their inner part with remaining grain boundary carbides and it is difficult to be precise. But new acicular Cr 7 C 3 carbides also appear among the initial TaC secondary carbides for Alloy2, and among the initial secondary TaC and Cr 7 C 3 for Alloy3. For the two alloys it is a consequence of an increase of the total carbon contained in the matrix alone. This clearly shows that the mechanism of integration of carbon into matrix also acts here. Like Alloy1 at 1100°C, it was seen for Alloy2 that a raising of the temperature unable this phenomenon which is probably also replaced by a disappearance of carbon from the alloy (becoming oxidised into gases). But this difference of behaviour remains to explain. Finally thermodynamic modelling allowed us to verify that these new compositions induced by oxidation are not prejudicial for the alloys sub-surface refractoriness.

Conclusion

It can be underlined that thermodynamic modelling, associated to microstructure phase quantitative analysis and microanalysis, was of a good help to better characterise and understand what it happened to our specific alloys and superalloys. This method allowed to supply the lack of microanalysis techniques for what quantitative measurements of contents for light and rare elements is a real problem. With this particular procedure it became possible to localise carbon, which is an element of a great importance in our alloys. Consequently curious phenomena have been identified.

Fig. 1 :

 1 Fig. 1: Example of deterioration of the sub-surface microstructure by oxidation

Fig. 2 :Fig. 3 :

 23 Fig. 2: Sub-surface microstructural changes for Alloy1 and correspondence between the different zones and the Cr profile

Fig. 4 :

 4 Fig. 4: Microstructural evolutions along the oxidation front of the two cobalt base alloys after oxidation at 1000°C for 100h

  Cr 7 C 3 and 1.0%surf TaC (surface fractions), then 2.6%mass of Cr 7 C 3 and 1.6%mass of TaC (mass fractions), for a local chemical composition of 27.7Cr-9.8Ni-2.1Ta (bal. Co) in weight percent. Alloy3 contains 10.1%surf Cr 7 C 3 and 0.7%surf TaC, then 8.0%mass of Cr 7 C 3 and 1.2%mass of TaC, for a local chemical composition measured of 28.9Cr-9.4Ni-1.6Ta (bal Co).

Fig. 6 : 1 ≅

 61 Fig. 6: Scheme of the carbon content evolution through the damaged zone of Alloy1 (1000°C)

Table 1 :

 1 Typical chemical compositions of the studied alloys (in weight %)

		Co	Ni	Cr	C	Ta
	Alloy1	/	bal	30	0.2	/
	Alloy2	bal	8.5	30	0.4	6
	Alloy3	bal	8.5	30	1.0	6

High temperature oxidation runs and microstructure examinations.

  Oxidation tests were performed at 1000°C, 1100°C and 1200°C for 100 hours, using a Setaram thermobalance under a flow of industrial air (80%N 2 -20%O 2 ) of about 1,5L/min. Samples were parallelepipeds, the dimensions of which were approximately 10mm x 10mm x 2mm. After tests, oxidised samples were covered with a gold layer by cathodic plasma pulverisation which allowed the external surface

	external
	Cr 2 O 3
	internal
	Ta 2 O 5
	lost of
	secondary and
	primary
	carbides
	precipitation of
	new phases (in
	dark)
	Co base superalloy oxidised at 1100°C for 100h

Table 2 :

 2 Values of carbon content limits and of resulting local solidus temperatures

		Locations	external limit	middle	internal limit acicular Cr 7 C 3	bulk
			of the carbideless zone	carbide zone
	Alloy1	T solidus (C=0)	1417	/	1406	1317	1317
	1000°C	T solidus (C=C max )	1381	/	1380
	Alloy1	T solidus (C=0)	1408	/	1393	/	1317
	1100°C	T solidus (C=C max )	1346	/	1350	/
	Alloy2	T solidus (C=0)	1455	/	1417	1277	1281
	1000°C	T solidus (C=C max )	1418	/	1388	/
	Alloy2	T solidus (C=0)	1446	1428	1392	/	1281
	1200°C	T solidus (C=C max )	1328	1323	1345	/
	Alloy3	T solidus (C=0)	1423	/	1393	1278	1277
	1000°C	T solidus (C=C max )	1391	/	1382	/


P Berthod

S Raude

A S Renck

C Rapin

R Podor

P Steinmetz

Influence of the Microstructural Texture of Cast Superalloys on their High-Temperature Oxidation Behaviour

Keywords: High Temperature Oxidation, Cast Superalloys, Grain Size, Dendritic Orientation

The high-temperature oxidation behaviours of Ni base and Co base cast superalloys were studied to determine the effect of different grain sizes and different surface dendritic orientations. These microstructural characteristics were obtained by varying solidification rate and cutting orientation with regard to the external surface. Thermogravimetry tests were run at 1000 and 1100°C and parabolic oxidation constants were considered. It appears that oxidation is faster for fine microstructures than for coarser ones for Ni or Co alloys including tungsten in their chemical composition, while it is the contrary when W is not present. When the sample surface is mainly parallel to the dendritic network, the oxidation rate is greater than for a surface perpendicular to dendrites, for the studied Ni alloy. The same phenomenon is observed for the studied Co base alloy at 1100°C but the order is inverted at 1000°C. These different behaviours can be explained by the grain boundary densities and orientations obtained on surface, since they can influence the diffusion of species involved in the oxidation phenomena. It is also possible that the characteristics of the chromia scale, such as grain size and general quality, depend on the microstructural texture of the alloy.

Introduction

The microstructure of metallic alloys resulting from classical foundry practice depends on their chemical composition but also on the solidification conditions. Indeed, local cooling rate and thermal gradient existing during solidification govern, respectively, grain size and dendritic orientation, as well as density and general orientation of grain boundaries. Knowing that the latter act as special diffusion paths for species like Cr atoms or O 2-anions [START_REF] Kofstad | High Temperature Corrosion[END_REF], it can be thought that the kinetics of high-temperature oxidation of superalloys could depend on both microstructural fineness and orientation. These points were studied for several Ni-base and Co-base cast superalloys which are strengthened by intergranular carbides of different types.

Effects of solidification conditions on microstructures

A metallic cast piece obtained by classical foundry practice usually presents a microstructural evolution from the external surface towards the last point to solidify [START_REF] Kurz | Fundamentals of solidification[END_REF][START_REF] Lesoult | [END_REF]4]. In the nearest neighbourhood of the mould wall, there is a thin zone where solidification begins by the appearance of numerous very small crystals (Fig. 1). Among the latter, only the ones which are the more favourably oriented with regard to the local thermal gradient (generally perpendicular to the mould wall), pursue their growth in the next zone. Thus, after this first zone is formed with a very fine and equiaxed microstructure, a columnar zone develops which displays an oriented microstructure. Here columns, or dendrites for Ni and Co base alloys, are mainly perpendicular to the mould wall, then parallel to the local thermal gradient existing during solidification. The third zone, where solidification finished, contains equiaxed dendrites that have grown in several directions, because of the absence of any very marked gradient. Whatever the zone considered, final crystal or dendrite sizes depend on the solidification rate or on the cooling rate. A low solidification rate leads to a coarser microstructure than does a high one. After machining, a cast piece can present locally external surfaces with different grain sizes and different orientations with regard to the sub-surface dendritic network.

Materials and methods

General experimental details. Four superalloys were synthesised using an induction furnace under argon gas, and solidified in sand moulds to give cylindrical pieces with a diameter of nearly 30mm. Chemical compositions of these alloys, a Ni base one and three Co base ones, are given in Table 1.

All alloys contained, in the grain boundaries, one or two types of carbides (eutectic M 23 C 6 and/or MC) known to be very useful for strengthening at high temperatures [START_REF] Bradley | Superalloys : A technical guide[END_REF][START_REF] Sims | The superalloys[END_REF]. Other experimental details (grain size). The following procedures were applied to all the alloys to allow the influence of grain size on the oxidation kinetics to be considered. For each alloy a sample for oxidation tests, with a coarse microstructure (characterized by a mean spacing between two secondary dendrite arms λ of about 50µm), was cut from the equiaxed zone of the castings. Elsewhere a part of the latter (about 50g) was remelted in a water-cooled copper crucible in another high induction furnace, in which the alloy solidified with a significantly higher cooling rate. Thus the same alloys were obtained with a finer microstructure (λ ≅ 25µm). An example of an alloy with a coarse and a fine microstructure is given in Fig. 2. the oxidation parabolic constants for S // (Kp // ) and S ⊥ (Kp ⊥ ) appear as the two unknown values of a linear equation (Eq. 4) in which also the measured total constant appears.

( )

+ + + = + ∆ + ∆ = ∆ (3) Kp . 2 Kp . 2 S S Kp . 2 S S // // = + ⊥ ⊥ (4)
Solving the two linear equations (Eq. 4) written for two samples of the same alloy oxidized at the same temperature, leads to values of both Kp // and Kp ⊥ .

Results

Effects of alloy grain size. Thermogravimetry results show that high-temperature oxidation kinetics are obviously dependant on the microstructural fineness. Indeed, for the same alloy, the mass gain curves obtained for the coarse and fine microstructures are different, as shown by Fig. 4. Table 2 presents all the oxidation parabolic constants determined from the obtained curves. These values also show that oxidation rate varies with grain size but not always in the same way. Differences between "coarse" and "fine" were very marked for 3 over the 6 cases, since the parabolic constants ratio were higher than 2 or lower than 0.5. The difference depended on the alloy, possibly on its chemical composition. For instance, the alloys showing a higher oxidation rate for a fine structure than for a coarser one all contained tungsten (Nickel1 and Cobalt1), while the opposite was found for the others (Cobalt2 and Cobalt3). For the only alloy tested for several temperatures (Cobalt2), the parabolic constants ratio increased with temperature, from values lower than 1 to values higher than 1. Thus, a fine Cobalt1 oxidized faster than a coarse Cobalt1 at relatively low temperatures (<1000°C), as fast as the coarse one at 1100°C, and slower than the coarse one at higher temperatures (>1200°C). Oxidation and/or diffusion phenomena seem to be a little different, since activation energies (Fig. 5) were 185kJ/mol for the coarse microstructure and 253kJ/mol for the fine one. The three points in Fig. 5 were not really aligned, doubtless because of a high volatilization rate of chromia at 1200°C (oxidised into gaseous CrO 3 ) which reduces the measured parabolic constant. If only the two points corresponding to the two lowest temperatures are considered, the activation energies become higher (Table 2). Another interesting observation is that a dependence of the oxidation rate on the density of carbides previously observed for this alloys family [START_REF] Martino | [END_REF], was observed here for the coarse microstructures but not for fine microstructures. Indeed, Kp coarse at 1000°C raised with C content, whatever the nature of the carbides, while all parabolic constants at 1000°C were similar when microstructures were finer (all values are near 1.5 10 12 . g 2 .cm -4 .s -1 ). and activation energies for the two grain sizes (Cobalt2 alloy) Fig. 5: ln(Kp) versus the inverse of temperature in Kelvin for the Cobalt2 alloy; determination of the activation energy associated to the coarse and the fine microstructures Microscope examinations were performed on the samples after special metallographical preparation. Even if the external Cr 2 O 3 scales were always mainly lost during cooling after thermogravimetry tests, some parts of this layer remained on the sample surfaces so that it was possible to estimate the chromia thickness. The order of the chromia thickness on the alloys was in most cases the same as for the parabolic constants. The depth of internal oxidation (Cr 2 O 3 and/or CrTaO 4 depending on the alloy) was also assessed and it appeared that there was no differences between the two microstructural finenesses for the same alloy (20µm for all alloys at 1000°C). Cobalt2 was also tested at two higher temperatures, 1100 and 1200°C. The order of the parabolic constant values and of the assessed chromia thicknesses were the same as at the lower temperature. Internal oxidation increased with temperature, from about 20µm (coarse or fine) at 1000°C to 30 (fine) or 50µm (coarse) at 1200°C. A zone having lost its initial intergranular carbides began to be visible at 1100°C and became deeper from the external surface at 1200°C, about 80 (coarse) or 90µm (fine), as illustrated by Fig. 6. Thus, with increasing temperature, oxidation rate and the carbide-free band growth rate increased and were higher for a fine microstructure than for a coarse one, while the internal oxidation depth coarse fine also increased but with an inverted order between coarse and fine. As was observed earlier [START_REF] Martino | [END_REF], this tendency remained at 1300°C since the depth of the carbide-free zone was, after 20 hours, nearly 120µm for a coarse microstructure and the double for a fine one. At these high temperatures, tantalum diffuses towards the oxidation front from the zone containing tantalum carbides. Tantalum oxides formed a discontinuous band in the alloy just under the chromia layer. Effects of dendritic orientation. The mass gain curves suggest that oxidation rate was more or less dependant on the dendritic orientation. Fig. 7 displays curves for Nickel1 at 1000°C which are sensibly different from one another, the oxidation rate being higher for the sample parallel to dendrites than for the perpendicular one. On the contrary, the kinetics were almost the same for Cobalt1 at 1000°C, with a tendency for a lightly slower oxidation rate for the sample mainly parallel to dendrites than for the perpendicular one. These observations are reflected in the parabolic constants values in Table 3, and amplified after uncoupling of the two contributions. Comparing the behaviour at 1000 and 1100°C, it appears that oxidation rate of Nickel1 was higher for samples mainly parallel to dendrites than for perpendicular ones, with a ratio of 10 at 1000°C and of 4 at 1100°C on the parabolic constants, and a ratio of 3 at 1000°C and of 2 at 1100°C for the rate of mass gain at each time. Similar results were reported earlier for simple Ni-30Cr-0.8C (%wt) alloys [START_REF] Berthod | poster presented at Gordon Research Conference "High Temperature Corrosion[END_REF] but with less marked differences. For Cobalt1, oxidation rate was lightly higher for the sample mainly perpendicular to dendrites than for the parallel one at 1000°C, and the difference increased with temperature. The external chromia thicknesses were of the same order as the uncoupled parabolic constants for Nickel1, with chromia scales thicker for the parallel surfaces than for the perpendicular ones. On the contrary there was no real difference of chromia scale thicknesses between the two orientations for Cobalt1. For the two alloys the depth of internal oxidation increased with temperature (it doubled between 1000 and 1100°C) and internal oxidation was more extended for Cobalt1 than for Nickel1. For Nickel1 the internal oxidation depth did not depend on the dendritic orientation (less than 20µm for 1000°C and about 35µm for 1100°C). On the contrary the internal oxidation depth was greater for the perpendicular orientation (30 at 1000°C and 65µm at 1100°C) than for the parallel one (20 at 1000°C and 40µm at 1100°C). The differences in oxidation rate observed for the two orientations appeared to be directly associated to the chromia growth rate in the case of Nickel1, while it was more related to the internal oxidation for Cobalt1.

Fig. 8: View of the two orientations of dendrites with regard to the surface (Cobalt1) and surface state after oxidation at 1100°C for 100 hours

Discussion

These results, taken together with earlier data [START_REF] Berthod | [END_REF], demonstrate that both grain size and dendritic orientation can affect the high-temperature oxidation rate. This influence remains to be fully explained ; however, a first hypothesis is presented here. The effect of alloy grain size on oxidation rate exists, but there is also a dependance on chemical composition, more precisely, on the presence of tungsten in the alloy. A fine microstructure leads to a high density of grain boundaries emerging at the surface. In that case, chromium can diffuse more rapidly towards the oxidation front, and this facilitates the rapid initial formation of the protective chromia scale and an efficient supply thereafter. Thus the chromia scale can be of a higher quality, and more protective. However, if the alloy contains tungsten, which participates in the formation of interdendritic carbides, oxidation of these leads to the formation of WO 3 gaseous oxides which can disturb the chromia growth and its protective quality. Thus, a small grain size promotes a high density of grain boundaries, which can lead to more severe deterioration of the chromia scale. Elsewhere, it is possible that the fine grain size of the alloy promotes the development of an external chromia layer with a high density of grain boundaries in the oxide itself. Its conductivity for ions or content of point defects may be higher, which may increase the oxidation rate. Dendrite orientation can have similar effects on the characteristics of the external chromia scale, and therefore on its protection efficiency. Since grain boundaries are often more conductive than the alloy matrix, it is likely that diffusion can be easier for the perpendicular orientation than for the parallel one, with consequences on the formation of the protective chromia scale and supply of Cr (Nickel1), and also on the internal oxidation progress (Cobalt1).

Conclusion

The chemical composition is not the only alloy parameter having some influence of the hightemperature oxidation behaviour of carbide-strengthened cast superalloys. Indeed, it appeared here that both grain size and dendritic orientation must be taken into account, for example, when comparative oxidation tests are made in order to optimise the chemical composition of an alloy for improving its oxidation behaviour. But this influence, which seems also to depend on other chemical or microstructural parameters, is difficult to define and more work remains to be done for better understanding of the involved phenomena. First, tests at other temperatures are needed on all alloys in order to define the activation energies. Second, microprobe contents profiles and texture analysis of the chromia layer are required to understand the diffusion phenomena and their differences when grain size or dendritic orientation vary.

Fig. 1 :

 1 Fig. 1: The three different microstructure zones usually existing from the mould wall in a metallic piece obtained by classical foundry practice

Fig. 2 :

 2 Fig. 2: Example of the two microstructural finenesses in the case of the Cobalt2 alloy

Fig. 3 :

 3 Fig. 3: Cutting modes for samples mainly parallel or mainly perpendicular to dendrites

Fig. 4 :

 4 Fig. 4: Some examples of mass gain curves for two different grain sizes of the same alloy (on the left: case of Nickel1 at 1000°C, on the right: case of Cobalt1 at 1200°C)

Fig. 6 :

 6 Fig. 6: Depth of the carbideless zone for Cobalt2 oxidized at 1200°C for 100 hours in the cases of a coarse microstructure (on the left) and of a fine one (on the right)

Fig. 7 :

 7 Fig. 7: Some examples of mass gain curves obtained for the two dendritic orientations

Table 1 :

 1 Chemical composition of the four cast superalloys studied

	weight % Co	Ni	Cr	C	W	Ta	Fe
	Cobalt1	bal.	9	30	0,4	6	3	/
	Cobalt2	bal.	9	30	0,4	/	6	/
	Cobalt3	bal.	9	30	1,0	/	6	/
	Nickel1	/	bal.	30	0,7	7,5	0	7,5
	Oxidation tests were performed at 1000°C and 1100°C for 100 hours, using a Setaram TGA92
	apparatus and under a flow of industrial air (80%N 2 -20%O 2 ) of about 1,5L/min. Samples were
	parallelepipeds, the dimensions of which were approximately 10mm x 10mm x 2mm. After testing,
	oxidized samples were covered by a gold layer by cathodic plasma pulverisation, which allowed the
	external surface to become electrically conductive. Then a sufficiently thick electrolytic nickel
	coating was deposed in order to mechanically protect the external chromia layer. Thereafter samples
	were cut, embedded in resin, polished and examined by Scanning Electron Microscope (SEM).

Table 2 :

 2 Values of the oxidation parabolic constants obtained for all tests

	10 12 . Kp	microstructure	Kp	fine	Activation	Activation
	g 2 .cm -4 .s -1	coarse Kp coarse	fine Kp fine	coarse ratio Kp	energy (kJ.mol-1)	energy (kJ.mol-1)
	Nickel1* (1000°C) Cobalt1* (1000°C)	0.57 1.15	1.55 1.44	2.72 1.25	for the 3 temperatures	for the 2 lowest temperatures
	Cobalt2 (1000°C)	2.69	1.52	0.53	coarse	fine	coarse	fine
	Cobalt 2 (1100°C)	18.7	15.3	0.82	185	253	282	336
	Cobalt2 (1200°C)	28.0	37.8	1.35
	Cobalt3 (1000°C)	3.46	1.39	0.41	* contains tungsten

values of 10 12 .Kp in g 2 .cm -4 .s -1

	1000°C	Kp plaq//	Kp //	1100°C	Kp plaq//	Kp //
		Kp plaq⊥	Kp ⊥		Kp plaq⊥	Kp ⊥
	Nickel1	1,71	2,30	Nickel1	5,57	6,28
		0,63	0,19		2,67	1,51
	Cobalt1	1,32	1,22	Cobalt1	7,14	5,76
		1,69	1,90		10,38	14,32

Table 3 :

 3 Values of the global parabolic constants and of the uncoupled ones for the two alloys


Patrice Berthod
email: patrice.berthod@centraliens-lille.org

Yassin Hamini

Lionel Aranda

Influence of tantalum on the rates of high temperature oxidation and chromia volatilization for cast (Fe and/or Ni)-30Cr-0.4C alloys

Keywords: Cast alloys, High temperature oxidation, Chromia volatilization, Tantalum carbides

Four cast superalloys, Fe-base and (Fe,Ni)-base alloys, all containing 30%Cr and 0.4%C, were elaborated with addition of 3% and 6% of tantalum. Their oxidation behaviours were studied at 1,000, 1,100 and 1,200°C during 50 hours. The oxidized surfaces of the samples were quantitatively characterized by measurements of thicknesses and surface fractions of internal oxides. The thermogravimetry files were treated according to the {m×(dm/dt) = Kp -Kv×m} equation, to obtain simultaneously the parabolic constant and the chromia volatilization constant. The internal tantalum oxides are more present in the Fe-base alloys and the carbide-free zones are less developed for the (Fe,Ni)-base alloys than for the others. The Fe-base alloys oxidize faster than the (Fe,Ni)-base and Ni-base alloys. The comparison with the corresponding Ta-free ternary alloys shows that the presence of Ta tends to accelerate the oxidation.

Introduction

The equiaxed superalloys obtained by classical foundry process are often reinforced by the presence of both interdendritic carbides appeared during solidification, and eventually also by secondary carbides obtained by applying a specific heat treatment. It is notably the case of cobalt-base alloys, but this principle of strengthening can also be used for nickel and iron-base alloys. Thus, in addition to chromium, the role of which is to bring a good resistance against high temperature oxidation to the conventionally cast superalloys [START_REF] Kofstad | High Temperature Corrosion[END_REF], tantalum, which is also a very oxidable element, can play a role in the high temperature oxidation of (Fe and/or Ni) -base alloys.

In this work the high temperature oxidation behaviours of two iron-base alloys and two ironnickel base alloys, containing 0.4%C and two amounts of tantalum 3% and 6%, are studied, with comparison with results obtained previously for ternary M30Cr0.4C alloys [START_REF] Berthod | to be published[END_REF] and two nickel-base alloys with 0.4C and 3% to 6%Ta [3].

Experimental Procedures

Elaboration of the alloys and thermogravimetry runs. Two Fe-base and two (Fe,Ni)-base alloys, containing 3 and 6 wt.% of tantalum (Table 1), were melt by HF induction and solidified in a watercooled copper crucible under pure argon (300mbar), to obtain ingots of about 100g from graphite and pure Fe, Ni, Cr and Ta (> 99.9% w , Alfa Aesar). The geometry of the samples for thermogravimetry experiments was about 2x10x10mm 3 . All samples were polished with SiC paper up to 1,200 grid. A Setaram TGA92 thermobalance was used to perform the thermogravimetric analyses, at 1,000°C, 1,100°C and 1,200°C during 50 hours under a dry industrial air flow at 1.5 L/h. The heating was realized at 20°C/min and the cooling after test at -10°C/min.

Metallography characterization.

After the thermogravimetry runs, the oxidized samples were coated by a cathodic evaporated gold film in order to allow the electrolytic deposition of nickel on their external oxide scale (when it remained on surface after cooling), to protect the latter against stresses due to cutting. The samples were cut and embedded in a cold resin, then polished with SiC paper from 80 to 1,200 grid under water, then finished with a diamond paste 6µm. The samples were examined using a Scanning Electron Microscope (SEM) Philips XL30. Generally it was done in the Back Scattered Electrons mode (BSE) and by using an accelerating voltage of 20kV. The Energy Dispersion Spectrometry device (EDS) of the SEM allowed analyzing the chemical composition of each sample in three different locations. This permitted the calculation of an average content for each element with its corresponding standard deviation value (results displayed in Table 1). The thickness of the chromia scales (or of the parts of this scale remaining on surface after cooling) and the depth of the carbide-free zones were measured in three locations and the average value was calculated. The internal oxidation was characterized by measuring, on three SEM micrographs of the sub-surface, the surface fraction of the sub-cortical oxides involving tantalum. These ones appear as white particles when they are observed in BSE mode. The average value of an equivalent thickness for this oxide was calculated according to Eq. 1 in which l phot is the length of the area covered by the micrograph and f S is the surface fraction of the white tantalum oxides determined using the Photoshop CS software.

s phot oxide ) Ta , Cr ( f l e × = . (1)
Kinetic constants of oxidation. The mass gain results must be first corrected from the progressive loss of carbon from the external surface which was oxidized into gaseous species. Indeed, thermodynamic calculations can show that the carbide-free zone which develops from the oxidation front cannot contain carbon. Correction was performed by adding, to the recorded files, the term R C (t) expressed in Eq. 2 representing the mass of lost carbon versus time (following a law supposed to be parabolic). In Eq. 2, ρ alloy is the density of the alloy, S the external surface of the sample, d cfz the final depth of the carbide-free zone, % w C alloy the carbon weight content of the alloy, t the time spent after the beginning of the isothermal stage and ∆t stage the total duration of this stage. (2)

The first oxidation constant K l , which characterizes the transient oxidation, corresponds to the slope of the straight line obtained at the beginning of the isothermal mass gain curves before the beginning of the actual parabolic oxidation. The parabolic constant Kp and the constant of chromia volatilization Kv were determined by using the method defined in an earlier work [4]. This was done by considering that the Kp and Kv constants are respectively equal to the ordinate at the origin and to the slope of the obtained straight line which appears when the thermogravimetry results are plotted according to Eq. 3 (in which m is the mass gain per unit area). The validity of these values was thereafter verified by comparing the experimental mass gain curve and the mathematical curve plotted using the obtained constants.

m Kv Kp dt dm m × - = × . ( 3
)

Results

Chemical compositions and microstructures. Table 1 presents the obtained compositions of the four studied alloys, measured by EDS analysis on three x250 areas. By considering the standard deviations, the targeted compositions appear to be well respected. The as-cast alloys present microstructures which are composed of a dendritic matrix and interdendritic carbides. The latter are exclusively tantalum carbides (TaC). These Fe-base and (Fe,Ni)-base alloys contain much more tantalum carbides than the corresponding Ni-base alloys with similar compositions [3] which contained both chromium and tantalum carbides. Surface microstructural states of the alloys after oxidation. The sub-surface microstructures of the oxidized alloys are illustrated by Fig. 1 (Fe-base alloys) and Fig. 2 (FeNi-base alloys), for the two extreme temperatures. The tantalum carbides which were initially present before test, disappeared from the oxidation front over a depth which increases with the temperature of the thermogravimetric test. The part of the alloys that is just under the external surface is affected by an internal oxidation that leads to dispersed "islands" of oxides of both chromium and tantalum. These ones seem to be denser and more developed for the alloys with 6%Ta than for the ones with only 3%Ta, and for 1,200°C and 1,100°C than for 1,000°C. There is also an external oxide scale, which remains sometimes adherent. This layer is mainly constituted of chromia Cr 2 O 3 but it can also contain films or particles of tantalum oxides. The thickness of the chromia and of the internal tantalum oxides are shown in Table 2. Values obtained for a Ni30Cr0.4C3Ta and a Ni30Cr0.4C6Ta alloys [3], oxidized during 100 hours (and not 50 hours), are also added for a comparison purpose. They are supposed to be close to the oxide thicknesses that were measured in this previous work but after multiplication by the square root of (50 hours/100 hours = 2). Indeed the growths of external oxides and also of internal oxides often fulfil a parabolic law.The corresponding values obtained for the ternary alloys Fe30Cr0.4C and Ni30Cr0.4C [START_REF] Berthod | to be published[END_REF] are added in the "0%Ta" column. (italic characters: the oxidation time for Ni30Cr0.4C3or6Ta was of 100h [3] and not 50 hours) (n. o. r. : no external oxide remained on the sample after cooling)

It appears that the chromia scale is thicker for the 3%Ta-containing alloys than for the ternary alloys. But the presence of 6%Ta involves a decrease of the chromia thickness which reaches again the values measured for the ternary alloys. Concerning the internal oxidation the equivalent thickness of tantalum oxide is more important for the 6%Ta-containing alloys than for the alloys with 3%Ta. This equivalent tantalum oxide thickness also decreases when Fe is progressively replaced by Ni for a same Ta content. The same observation can be done for the Cr 2 O 3 thickness. Table 3 displays the average depths of the carbide-free zones. Concerning the Fe-base alloys, the carbide-free zone is more extended for a tantalum-free alloy than for its corresponding 3%Tacontaining alloy, and the carbide-free zone depth goes on decreasing when the Ta content increases up to 6%. Concerning the effect of a change of the base elements, it appears that the depth of the carbide-free zone reaches a minimal value when the contents of Ni and Fe are similar.

Thermogravimetry results.

As examples, the mass gain curves obtained for the (Fe,Ni)-base alloys at all temperatures are presented in Fig. 3. A lot of curves are affected by one or several mass gain jumps due to local detachments of the external chromia layer. The Ta-containing alloys oxidize faster than the corresponding Ta-free alloys, while the 3%Ta-containing alloys and the 6%Tacontaining alloys have mass gain curves which are often very close to each other. An example of treatment of the mass gain results according to the {m × (dm/dt) = Kp -Kv × m} equation is given in the graph of Fig. 4. The values obtained for the two kinetic constants were tested by verifying that the experimental curve and the mathematical curve plotted with the values obtained for Kp and Kv are or not superimposed. Generally the constants were effectively validated.

The values obtained for the linear K l constant are given in Table 4, and the values of the parabolic constant Kp and of the volatilization constant Kv are presented in Table 5. Very high values of both Kp and Kv were found for all the Fe-base alloys for 1,200°C, while they are on the contrary on the same order of magnitude as the (Fe,Ni)-base and Ni-base alloys for the two lower temperatures. For the parabolic constant of these Fe-base alloys at all temperatures, one can notice an increase of its value between 0% and 3%Ta, then decreasing between 3%Ta and 6%Ta. This tendency globally remains for the (Fe,Ni)-base alloys and for the Ni-base ones, except for 1,000°C for which the parabolic constant does not vary with the Ta content for the (Fe,Ni)-base alloys and decreases for the Ni-base alloys. Concerning the base elements, it can be noted that, almost at all temperatures, the Kp constant decreases when Fe is progressively replaced by Ni.

Concerning the Kv constant, there are globally an increase of its value when more Ta is present in the alloy (whatever the base element) and a decrease when Fe is replaced by Ni as base element. Discussion. In these Fe-base and (Fe,Ni)-base alloys the presence of Ta leads to similar densities of tantalum carbides, as already seen in Co-base alloys with the same C and Ta contents, and significantly higher than for Ni-base alloys [3]. Tantalum seems diffusing from the carbides towards the oxidation front to be oxidized in internal oxides, a phenomenon that promotes the development of a carbide-free zone. The depth of the latter is higher for a higher temperature of oxidation, but also higher for a lower Ta-content in the alloy. Thus, even if the equivalent thickness of tantalum oxide is generally slightly higher for a Fe-base alloy with a double Ta content, oxidation seems to involve almost the same quantity of tantalum, then promotes, in the Fe-base alloy with 3%Ta, the disappearance of tantalum carbides over a double depth compared to the Fe-base alloy with 6%Ta. The ratio of carbide-free depths between the (Fe,Ni)-base alloys with 3%Ta and 6%Ta, or between the Ni-base alloys with 3%Ta and 6%Ta, are not the same. A significant part of Ta is also contained in the matrix of these alloys while Ta is almost only present in the form of carbides in Fe-base alloys. The location of tantalum in the alloy seems to have consequences on the importance of internal oxidation, more precisely on the oxidation of tantalum. When this element is in grain boundaries as tantalum carbides (Fe-base alloys: higher equivalent thickness of tantalum oxide and deeper carbide-free zones) its diffusion towards the oxidation front seems easier than when a part of tantalum is in solid solution in the matrix ((Fe,Ni)-base and Ni-base alloys). This oxidation of tantalum can be the reason of the higher mass gains at 1,200 and 1,100°C of the alloys containing tantalum, compared to the ones for the ternary alloys. Oxidation of tantalum can also be responsible of the difficulties that were sometimes met for the (Kp+Kv) determinations and validations, since this second oxidation phenomenon, without volatilization of the formed oxide, is not taken into account in the {m × dm/dt = Kp -Kv × m} law [4]. Then, the simultaneous determination of both Kp and Kv is probably more accurate in the cases where tantalum is little involved in oxidation, i.e. when it is less present or when it is more difficult for it to diffuse. The latter condition is met for the lowest test temperatures, for the lowest tantalum contents, for Ni-base alloys rather than for (Fe,Ni)-base and Fe-base alloys, and for low carbon contents.

The faster oxidation of the Fe-base alloys compared to the (Fe,Ni)-base and Ni-base alloys can be related to the crystallographic structure of the matrix at the test temperatures, i.e. BCC (ferritic) for the Fe-base alloys and FCC (austenitic) for the others. The nature of the major element itself also plays a role since oxidation is also slightly faster for a (Fe,Ni)-base alloy than for its corresponding Ni-base alloy. The decrease of the parabolic constant of a Fe-base alloy when the Ta content is changed from 3 to 6%, is also different to what it is encountered for the austenitic alloys. There is probably an effect of the BCC-FCC difference and/or of the base element difference here too. The obtained Kv values are in accordance with data from literature since, for austenitic alloys at 1,100°C, they remain near 37 [5] and 83 [6] x 10 -10 g.cm -2 .s -1 , and at 1,200°C near 124 [5] and 280 [6] x 10 -10 g.cm -2 .s -1 . It is also often the case for the ferritic alloys, excepted for 1,200°C for which the Kv values which were determined are very high. Bold characters: excellent fit between the experimental mass gain curve and the mathematical one Normal characters: good fit between the experimental mass gain curve and the mathematical one Italic characters: not good fit between the experimental mass gain curve and the mathematical one

Conclusion

When present in Fe-base and (Fe,Ni)-base carbon-containing alloys, tantalum leads to a TaC carbides network which seems to influence the high temperature oxidation behaviour. Their mass gains are higher than for Ta-free alloys, especially for the highest temperatures. Moreover the oxidation of Ta leads to difficulties for determining the parabolic and chromia volatilization constants by using the {m × dm/dt = Kp -Kv × m}-treatment. This is particularly true in the case of iron-base alloys in which the diffusion of Ta towards the oxidation front is the easiest because of its availability in grain boundaries. Thus, an efficient reinforcement by tantalum carbides is probably easier to obtain in Fe-base and (Fe,Ni)-base alloys than in Ni-base alloys, but unfortunately with a significant decrease of the resistance against high temperature oxidation of alloys.

  Fig. 1. Cross section metallographies of the two Fe-base alloys containing Ta after oxidation for 50 hours at the two extreme temperatures 1,200°C (top) and 1,000°C (bottom) FeNi30Cr0.4C3Ta FeNi30Cr0.4C6Ta Fig. 2. Cross section metallographies of the two (Fe,Ni)-base alloys containing Ta after oxidation for 50 hours at 1,200°C (top) and 1,000°C (bottom)

Fig. 3 .

 3 Fig. 3. Mass gain curves of the FeNi30Cr0.4CyTa (y = 3 and 6) alloys at 1,200 (a), 1,100 (b) and 1,000°C (c)

  Fig. 4. {m × (dm/dt) = Kp -Kv × m}-type treatment of the curve for FeNi30Cr0.4C6Ta at 1,100°C

Table 1 .

 1 The obtained chemical compositions for the alloys (average values ± standard deviation)

	Weight contents [%]	Fe	Ni	Cr	Ta	C
	Fe30Cr0.4C3Ta	Bal.	/	29.5 ± 2.7	2.4 ± 0.4	≈ 0.4
	Fe30Cr0.4C6Ta	Bal.	/	29.2 ± 1.6	6.2 ± 1.1	≈ 0.4
	FeNi30Cr0.4C3Ta	Bal.	34.0 ± 0.4	30.0 ± 0.3	2.8 ± 0.4	≈ 0.4
	FeNi30Cr0.4C6Ta	Bal.	32.5 ± 0.2	29.5 ± 0.3	6.4 ± 0.3	≈ 0.4

Table 2 .

 2 Values of the average thicknesses of external chromia and of the equivalent (Cr,Ta) oxide assessed by image analysis (in each case average of 3 measures +/-1 std deviation)

	Av. thick.	T	0%Ta		3%Ta		6%Ta
	of Cr oxide Equ. thick. of Ta oxide	(°C)	Cr 2 O 3 (µm)	Cr 2 O 3 (µm)	(Cr,Ta) ox. (µm)		Cr 2 O 3 (µm)	(Cr,Ta) ox. (µm)
		1200 18 ± 3	32 ± 6	7.4 ± 0.4		16 ± 4	8.0 ± 2.0
	Fe30Cr0.4C	1100 18 ± 3	22 ± 2	5.3 ± 1.0		19 ± 3	7.3 ± 0.8
		1000 6.6 ± 0.7	11 ± 1	0.7 ± 0.2		8.3 ± 1	2.2 ± 0.3
		1200	/	30 ± 2	4.8 ± 0.3		31 ± 2	9.2 ± 1.1
	FeNi30Cr0.4C	1100	/	13 ± 2	1.8 ± 0.3		12 ± 2	4.4 ± 0.9
		1000	/	10 ± 1	0.7 ± 0.3		6.2 ± 2	1.1 ± 0.4
		1200 21 ± 3	n. o. r.	5.9 ± 0.7 (4.2*)	n. o. r.	4.8 ± 2.3 (3.4*)
	Ni30Cr0.4C	1100 11 ± 1	13 h 100 / h 50	, i.e. divided by 2 )

± 4 (9.2*) 2.9 ± 0.7 (2.1*) 8.6 ± 1.0 (6.1*) 6.0 ± 0.6 (4.2*) 1000 4.6 ± 0.7 13 ± 4 (9.2*) 1.9 ± 0.2 (1.3*) 6.0 ± 1.0 (4.2*) 2.1 ± 0.2 (1.5*) (*: values deduced from the former multiplicated by

Table 3 .

 3 Values of the average depths of the carbide-free zones (in each case average of 3 measures +/-1 std deviation)

	Average depth of		0%Ta				3%Ta	6%Ta
	the carbide-free zone	Temperature	d cfz (µm)				d cfz (µm)	d cfz (µm)
		1200°C	333 ± 33				272 ± 11	91 ± 9
	Fe30Cr0.4C	1100°C	144 ± 17				98 ± 2	65 ± 4
		1000°C	41 ± 2				36 ± 6	20 ± 2
		1200°C	/				103 ± 10	107 ± 6
	FeNi30Cr0.4C	1100°C	/				49 ± 7	40 ± 4
		1000°C	/				19 ± 5	11 ± 1
		1200°C	143 ± 2			211 ± 10 (149*)	190 ± 8 (134*)
	Ni30Cr0.4C	1100°C	66 ± 5				79 ± 8 (56*)	76 ± 4 (54*)
		1000°C	26 ± 3				30 ± 2 (21*)	21 ± 4 (15*)
	(*: values deduced from the former multiplicated by	50	h	/	100	h	, i.e. divided by 2 )
	(italic characters: the oxidation time for Ni30Cr0.4C3or6Ta was of 100h [3] and not 50 hours)
	(n. o. r. : no external oxide remained on the sample after cooling)

Table 4 .

 4 Values of the transient linear K l constant (x 10 -8 g.cm -2 .s -1 )

	K l	temperature	0% Ta	3% Ta	6% Ta
		1200°C	75	128	88
	Fe30Cr0.4C	1100°C	29	29	48
		1000°C	4.9	4.7	15
		1200°C	/	60	49
	FeNi30Cr0.4C	1100°C	/	18	21
		1000°C	/	11	10
		1200°C	47	36	63
	Ni30Cr0.4C	1100°C	20	10.5	23
		1000°C	6.5	8.7	3.7

Table 5 .

 5 Values of Kp (x 10 -12 g 2 .cm -4 .s -1 ) and Kv (x 10 -10 g.cm -2 .s -1 ) constants

	Kp	temperature	0% Ta	3% Ta	6% Ta
		1200°C	499	844	696
	Fe30Cr0.4C	1100°C	59	131	119
		1000°C	6.3	87	13
		1200°C	/	213	204
	FeNi30Cr0.4C	1100°C	/	43	65
		1000°C	/	8.5	10
		1200°C	70	218	248
	Ni30Cr0.4C	1100°C	18	39	42
		1000°C	7.1	5.8	3.2
	Kv	temperature	0% Ta	3% Ta	6% Ta
		1200°C	514	436	770
	Fe30Cr0.4C	1100°C	73	98	137
		1000°C	27	348	-3.6
		1200°C	/	213	200
	FeNi30Cr0.4C	1100°C	/	53	117
		1000°C	/	51	23
		1200°C	82	156	235
	Ni30Cr0.4C	1100°C	58	26	70
		1000°C	32	16	24


Patrice Berthod
email: patrice.berthod@centraliens-lille.org

Pierric Lemoine

Lionel Aranda

Study of the behavior in oxidation at high temperature of Ni, Co and Fe-base alloys containing very high fractions of carbides

Keywords: Nickel alloys, Cobalt alloys, Iron alloys, High temperature oxidation, High carbon contents

Nine cast alloys reinforced by very high fractions of carbides, Ni-30Cr-xC, Co-30Cr-xC and Fe-30Cr-xC with x varying from 1.2 to 2.0, were tested in oxidation at high temperature between 1,000 and 1,200°C in air for 50 hours. After oxidation, their surfaces and sub-surfaces were characterized. Even for very high carbon contents, the chromia-forming behaviour of the nickel alloys is kept. The oxidation modes of the cobalt alloys and iron alloys are not changed compared to low carbon alloys of these families. The differences of diffusion easiness of chromium in matrix, between nickel alloys, cobalt alloys and iron alloys are the same as for alloys with lower carbon contents, as suggested by the lower chromium gradients in the nickel alloys compared to the two other alloy types. Sub-surface microstructure transformations due to oxidation were observed in some cases (coarsening of carbides due to an inwards diffusion of carbon, change of the sharing between BCC-FCC of iron matrix due to outwards diffusion of chromium). Catastrophic oxidation never occurred for these alloys during the 50 hours of exposition to air at high temperature.

Introduction

Some of the alloys or superalloys for uses at high temperature contain significant quantities of carbides, the role of which is essentially to achieve sufficient mechanical properties [START_REF] Bradley | Superalloys: A technical guide[END_REF][START_REF] Sims | The superalloys[END_REF]. In alloys which contain chromium to be resistant against high temperature corrosion [START_REF] Kofstad | High Temperature Corrosion[END_REF], some of the carbides may involve a part of chromium since this one is a carbide-former element. Thus, for a given chromium content, which is generally 30%wt or lower, high carbides fractions lead to a new distribution of chromium in alloy, with an especially great content in grain boundaries and therefore an impoverishment in matrix. This can be of great importance for the high temperature oxidation behaviour of the alloys.

In this work the deterioration, by oxidation in air at high temperature, of the surface and the subsurface of three families of alloys (Ni-base, Fe-base and Co-base), all containing 30%wt of chromium and several high contents of carbon, was studied by metallographic examinations after oxidation tests.

Experimental Procedures

Elaboration of the alloys and oxidation tests. Nine alloys M-30%Cr-x%C (x=1.2, 1.6 and 2.0), the base elements of which are the ones of most of the conventionally cast superalloys for high temperature applications, i.e. nickel, cobalt and iron, were elaborated by High Frequency induction melting (of graphite and pure Ni, Co, Fe and Cr, > 99.9%wt., Alfa Aesar) in an inert Argon atmosphere. Solidification occurred in the water-cooled copper crucible of the CELES induction furnace. Ingots of about 100g were obtained and samples for oxidation test were cut, with a parallelepipedic geometry of about 3 x 10 x 10mm 3 . After a polishing with SiC paper up to 1,200 grid, the samples were exposed to oxidation by air for 50 hours in a tubular furnace, at 1,000°C, 1,100°C or 1,200°C.

Metallographic characterization.

After oxidation, an electrolytic nickel layer was deposited on the samples, the oxidized surface of which was become electrically conductive with a first deposition of gold by cathodic evaporation. The samples were cut in two parts using an Accutom-2 saw, which were thereafter embedded in a cold resin (Araldite CY230 + Strengthener Escil HY956). Polishing was performed first with SiC paper from 80 to 1,200 grid under water, and second with a diamond paste 6µm (DP-Suspension P of Struers). The samples were examined using a Scanning Electron Microscope (SEM) Philips XL30, in the Back Scattered Electrons mode (BSE) under an accelerating voltage of 20kV. The thicknesses of the external oxide scales (or of the parts of this scale remaining on surface after cooling), and the depths of the carbide-free zones, were measured in three locations, with calculations of the average value and the standard deviation. Two profiles of concentrations were performed across the zone affected by oxidation, by using a Cameca SX50 microprobe in the Wavelength Dispersion Spectrometry mode, in order to assess the chemical composition changes in this zone and evaluate the chromium depletion.

Results

Chemical compositions and microstructures of the alloys. Table 1 presents the designation and the targeted composition of the nine studied alloys. The two last columns also give a description of the microstructure of the alloys at the intermediate temperature 1,100°C, calculated using Thermo-Calc [START_REF]Foundation for computational thermodynamics[END_REF] with a database containing the descriptions of all the M-Cr-C (M=Ni, Co or Fe) systems and their sub-systems. The targeted compositions, which were well respected by the contents really obtained, were used for performing the thermodynamic calculations. These ones showed that the bulk of the nickel-base alloys are constituted of a FCC matrix and Cr 7 C 3 carbides (stoichiometry verified by real WDS spot analysis) in the [1,000 -1,200°C] temperature range. The volume fraction of Cr 7 C 3 is about 15% at 1,100°C for the N12 alloy and increases up to almost 25% for the N20 alloy. The matrix of the cobalt alloys is also FCC between 1,000 and 1,200°C, and carbides are also Cr 7 C 3 (verified by WDS) with similar volume fractions as the nickel alloys. The microstructure of the iron alloys is more complicated since the nature of the matrix varies with the carbon content. This one is only BCC for F12, both BCC and FCC for F16, and essentially FCC for F20. In addition, the matrix nature also varies a little with temperature, even in the 1,000-1,200°C. These variations of the matrix nature are related to the chromium content in the matrix which depends on the fraction of chromium carbides, i.e. BCC when Cr is more present in matrix and FCC when it is less present because a higher part of this element is involved in the carbides. However, the nature of the present carbides is constant since they are Cr 23 C 6 in the three iron-base alloys for the three temperatures. Their volume fractions, which are higher than for the Cr 7 C 3 in the other alloys, also increases with the carbon content and they can become higher than 30% with the F20 alloy.

Surface states of the oxidized nickel-base alloys. The oxidized nickel alloys are all covered by an external oxide scale which is essentially chromia (Fig. 1). The thickness of this scale increases which temperature (from about 7µm for 1,000°C up to about 20µm for 1,200°C) but is seemingly not clearly dependent on the carbon content of the alloy. Just under this external scale, an outer zone of the alloy, in which the carbides have disappeared, developed from the external surface. The depth of this carbide-free zone increases with temperature, and also seems to depend on the carbides density. Indeed, it tends to be a little thinner for the N20 and N16 alloys than for N12. In the carbides-richest N20 alloy, in which carbides initially exist with the highest volume fraction and with a homogenous repartition, the carbides seem to be replaced by voids in the carbide-free zone (Fig. 5). One can also notice that there are carbides, in the sub-surface of the N12 alloy oxidized at 1,000 and 1,100°C, which seem to be coarsened in an alloy zone along the inner frontier of the carbide-free zone. This can result to an inwards diffusion of carbon released by the disappearing carbides [START_REF] Berthod | [END_REF]6]. The depth of the carbide-free zone was measured in three different locations and this led to both average values and standard deviations. They were plotted in the first graph of Fig. 4, in which previous results obtained for a Ni30Cr0.8C alloy oxidized at the same temperatures during the same time are added [7]. This graph illustrates the dependence of the carbide-free zone depth on both temperature and the initial carbides density in the alloy. But there are also some exceptions (e.g. position of the N16 oxidized at 1,100°C).

Surface states of the oxidized cobalt-base alloys. After oxidation, the cobalt alloys are also covered by an external oxide layer (Fig. 2), which is not so homogenous as for the nickel alloys. Both its nature and its thickness evaluate along the external surface. Concerning the nature, as showed by spot WDS microanalysis, some oxides are Cr 2 O 3 and others are CoO. The two forms are clearly separated from each other, which can be easily seen since they present two different levels of grey when observed with the SEM in BSE mode (Cr 2 O 3 darker than CoO). Like it was earlier observed for Co30Cr alloys containing less carbon (e.g. in [8]), chromia and cobalt oxide can be either superposed (inner CoO and outer Cr 2 O 3 ) or juxtaposed. In addition, oxidation locally penetrates in the alloy (e.g. C20 oxidized at 1,000°C), and in other areas the alloy itself seems to be replaced by oxides (especially for the two highest temperatures). At 1,200°C, there is also a significant attack of the interdendritic spaces in the carbide-free zone. There is other cases for which oxidation can be locally very fast and accompanied by an enhanced interdendritic attack (Fig. 5).

Concerning the depth of the carbide-free zone, an increase with temperature is still observed (second graph in Fig. 4). However, on the contrary with what was seen for the nickel-alloys, there is no dependence of the depth of the carbide-free zone on the density of the interdendritic carbides network, even if the results obtained for Co30Cr alloys containing less carbon are added [8]. For the two lowest temperatures, the carbide-free zones of the cobalt-base alloys are seemingly as deep as for the nickel-base alloys, except for the temperature of 1,200°C for which the carbide-free zone is clearly more extended for the cobalt alloys than for the nickel alloys.

Surface states of the oxidized iron-base alloys. The oxidized iron-base alloys are also covered by an external oxide scale, the thickness of which is generally variable along the extreme surface of the sample. The nature of this oxide also varies, since either two different oxides can be observed separately (superposed or juxtaposed), or a more continuous change of the Cr and Fe contents in the (Cr,Fe) 2 O 3 stoichiometry can be revealed across the external scale. A carbide-free zone developed inwards from the extreme surface here too, in which internal oxidation occurred, sometimes especially in the interdendritic spaces. The oxidation front is quite irregular, like for the cobalt alloys and differently to the nickel alloys. The depth of the carbide-free zone always increases with temperature and it is similar to the ones measured for the nickel alloys and for the cobalt alloys for 1,000 and 1,100°C. It is also similar to the nickel-base ones for 1,200°C, i.e. lower than for the cobalt alloys. The difference of carbide-free depth with the 0.8%wt.C-containing alloy studied in [8] is more marked than for the nickel alloys. It also be noticed that the sharing of the matrix of the F16 alloy between the BCC form and the FCC form, which already existed in the whole alloy before oxidation, seems to be modified in the inner part of the zone affected by oxidation. Indeed the pale areas are more frequent and extended than in the bulk (Fig. 5).

Chromium profiles across the zone affected by oxidation. Two chromium content profiles were performed by WDS microanalysis per oxidized sample, across the zone affected by oxidation (including the carbide-free zone, the zone in which carbides have changed when it exists, and in an outer part in the bulk). The results are given in Table 2, in which the average values (calculated with three measures for each sample) of the depth of the carbide-free zone is reminded for comparison.

Of course all samples display a chromium depletion from the external surface, which is more or less extended, this depending on the base element, the carbides fraction and the temperature). In most cases, the zone of Cr-depletion obviously corresponds to the carbide-free zone. But it is true that the depletion depth is also often more extended than the carbide-free zone. For the chromium depletion, one finds again the same phenomenon as already met above for the carbide-free depth: the chromium depletion depth increases when the temperature raises and when the carbon content or the carbide fraction decreases. Table 2 also presents the values of the average Cr content in the extreme surface for all samples. Generally these Cr contents on surface are higher for the nickel alloys than for the cobalt alloys, and for the latter than for the iron alloys. This Cr content on surface also decreases when the carbides fraction increases, and finally the Cr content on surface increases with temperature. The average values of the Cr gradient across the zone affected by oxidation are also given. All Cr gradients decrease when the temperature increases, and they are generally lower for the nickel alloys than for the cobalt alloys, and lower for the latter than for the iron alloys. Their dependence on the carbide fraction is clear in the case of the nickel alloys and the cobalt alloys since the gradient decreases when the carbide fraction increases, at all temperatures (but there are one or two exceptions). It is more complicated in the case of the iron alloys for which it was sometimes impossible to define a gradient.

Discussion

These (Ni, Co or Fe)-30Cr-1.2 to 2.0C alloys containing very high carbides fractions all displayed a high temperature oxidation behaviour which appeared to be never catastrophic during 50 hours at temperatures which were very high, although they contains high atomic contents of carbon. Initially one can think that this element either will induce itself a catastrophic behaviour, because of its oxidation into gaseous species, or will promote a lost of the chromia-forming behaviour of the emerging matrix. Concerning this second point, the chromium content in matrix was effectively very lowered by the precipitation of high fractions of chromium carbides in grain boundaries. In fact, the chromia-forming behaviour of the nickel alloys was kept, even when the temperature was as high as 1,200°C, while the cobalt alloys and the iron alloys, which are not really Cr 2 O 3 -forming at these temperatures when they contain less carbon [8], did not really change of behaviour. Indeed, the formation, on the external surface, of Cr 2 O 3 and CoO together, or of (Cr,Fe) 2 O 3 , was already encountered respectively for Co -30Cr -0 to 0.8%C alloys and Fe -30Cr -0 to 0.8%C alloys, in a previous work [8]. Here too, oxidation penetrated inwards the (cobalt or iron) alloy with an irregular front, since first the whole alloy began to be attacked, and second oxidation locally accelerated in several points of the oxidation front.

In all cases a carbide-free zone developed from the external surface because of the participation to the oxidation on the external surface, of the chromium contained in the carbides, together with chromium present in solid solution in matrix. For a same carbon content, the carbide-free zone is less extended for the iron alloys than for the nickel alloys, probably because the penetration of the oxidation front into the iron alloys since both Cr and Fe are obviously oxidized together. About the cobalt alloys, for which cobalt is also oxidized in the same time as chromium, the very deep carbide-less zone in some cases are probably related to a higher rate of oxidation than for nickel alloys. This was not easy to verify on the post mortem samples since oxides partly quitted the samples during cooling and the quantity of oxides observed by metallography was then minimized.

Chromium diffuses both along grain boundaries and through the matrix. The second phenomenon is more or less easy, this depending on the temperature, the nature of the matrix (chromium gradients in nickel matrix lower than in cobalt or iron matrixes) and, seemingly, with the density of the carbides network. Indeed, the Cr gradient appeared to be lowered by an increased carbon content. One can also notice the matrix change of the 1.6%wt.C-containing iron alloy, which is probably due to the small impoverishment of chromium in the inner part of the zone affected by oxidation. In this zone of this alloy, the carbon content of which led to an intermediate Cr content in matrix allowing the presence of the BCC and FCC forms simultaneously, a part of BCC (favoured by Cr) appeared to be effectively converted in FCC.

Conclusion

Even when they content high fraction of carbides, e.g. for achieving special mechanical properties, the Ni-30Cr based alloys, but also the Co-30Cr and Fe-30Cr based alloys, are able to resist oxidation at least for several tenths hours at high temperature. But such alloys are especially exposed to catastrophic oxidation (which never occurred here) because of the great carbon content and of the low Cr contents in matrix due to the precipitation of high fraction of chromium carbides.

N16 after 50h at 1 Fig. 1 .Fig. 2 .Fig. 3 .

 1123 Fig. 1.

Fig. 4 .Fig. 5 .

 45 Fig. 4. Evolution of the depth of the carbide-free zone with the carbon content and the temperature for the Ni-base (top), Co-base (middle) and Fe-base (bottom) alloys

Table 1 .

 1 Designations, compositions and microstructures of the studied alloys

			Targeted chemical composition (%wt.)		Microstructures at
	Designation						1,100°C (vol. frac. %)
		Ni	Co	Fe	Cr	C	Cr 23 C 6	Cr 7 C 3
	N20	Bal.	/	/	30	2.0	/	24
	N16	Bal.	/	/	30	1.6	/	19
	N12	Bal.	/	/	30	1.2	/	15
	C20	/	Bal.	/	30	2.0	/	23
	C16	/	Bal.	/	30	1.6	/	19
	C12	/	Bal.	/	30	1.2	/	14
	F20	/	/	Bal.	30	2.0	32	/
	F16	/	/	Bal.	30	1.6	28	/
	F12	/	/	Bal.	30	1.2	21	/

Table 2 .

 2 Comparison between the average values of the carbide-free depths and of the Cr-depleted depths (WDS microanalysis); average values of the Cr-content in extreme surface and of the Cr weight content gradient

		1,000°C during 50 hours	1,100°C during 50 hours	1,200°C during 50 hours
	Alloy	carb.-free d.	Cr deplet.	carb.-free d.	Cr deplet.	carb.-free d.	Cr deplet.
		(µm)	(µm)	(µm)	(µm)	(µm)	(µm)
	N20	12	30	27	30	56	30
	N16	13	15	44	50	61	50
	N12	17	25	30	30	84	70
	C20	20	20	42	40	106	100
	C16	9	25	49	40	114	100
	C12	19	25	40	75	115	200
	F20	13	10	30	35	40	?
	F16	14	20	36	60	58	70
	F12	21	20	40	50	59	90
		1,000°C during 50 hours	1,100°C during 50 hours	1,200°C during 50 hours
	Alloy	minimal	Cr gradient	minimal	Cr gradient	minimal	Cr gradient
		%wt. Cr	(%wt./µm)	% wt. Cr	(%wt./µm)	% wt. Cr	(%wt./µm)
	N20	14	0.058	14	0.019	18	0.016
	N16	17	0.069	16	0.060	18	0.012
	N12	18	0.164	20	0.057	20	0.021
	C20	13	0.234	7	0.351	16	0.047
	C16	12	0.461	15	0.138	16	0.062
	C12	18	0.525	10	0.180	17	0.074
	F20	13	?	10	0.518	?	?
	F16	11	0.727	11	0.251	18	0.057
	F12	15	0.368	11	0.238	15	0.100