Elodie Conrath
email: elodie.conrath@univ-lorraine.fr

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

Kinetics of High Temperature Oxidation and Chromia Volatilization for HfC-containing Nickel-based Alloys

Keywords: Nickel-based alloys, HfC carbides, High temperature oxidation, Kinetic constants, Scale characteristics

Many cast chromia-forming nickel-based superalloys are reinforced by carbides. In such alloys the primary chromium carbides or tantalum carbides rapidly lose their strengthening effect in service at high temperature because of decrease in volume fraction and morphology evolution. Other carbides, HfC for example, are more stable at high temperature and they can be candidates for the reinforcement of this type of superalloy. In this work, three nickel-based alloys containing 25 wt.%Cr, 0.25 or 0.50 wt.%C, and Hf with contents high enough (3.7 and 5.6 wt.%) to promote the formation of numerous primary HfC carbides, were elaborated by foundry. They were tested in oxidation by air during 46h at 1200°C, with thermogravimetry measurement and metallographic characterization. All the mass gain curves obtained are parabolic and the oxidation rates of the studied alloys are only slightly faster than for corresponding ternary nickel alloys containing the same chromium and carbon quantities. The obtained values of the parabolic constant k p and of the chromia volatilization constant k v , deduced by applying the {m × dm/dt = k p -k v × m} method, are typical of a chromia-forming behavior. However, small quantities of HfO 2 and NiCrTaO 4 oxides are noticed in addition to the chromia scale. In the bulk, the volume fraction and morphology of the HfC carbides only changed a little.

INTRODUCTION

The γ/γ' single crystal nickel-based superalloys are among the best refractory alloys for the applications at very high temperature [START_REF] Donachie | Superalloys : A Technical Guide[END_REF]. Their high contents in aluminum allow them resisting high temperature oxidation by isolating the metallic alloy from the aggressive atmosphere with a protective continuous alumina scale [START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF]. In some cases a good resistance against corrosion by liquid substances (molten glass for example) is also required. In such case the alloys must contain at least twenty weight percents of chromium, with as consequence a decrease in their aluminum content to low values. An efficient strengthening by high volume fractions of γ' precipitates cannot be achieved in these conditions and the alloy's reinforcement must be ensured by other particles, primary carbides for example. These ones are generally chromium carbides [START_REF] Vasseur | [END_REF] since, in cast chromium-rich nickel-based alloys, tantalum carbides are often much less stable than in cobalt-based alloys [4]. It was seen in a previous work that another type of MC carbides, hafnium mono-carbide, is thermally more stable than tantalum carbides when present in chromium-rich cobalt-based alloys [5]. In addition, in a more recent work, it was demonstrated that primary hafnium carbides can be successfully obtained in chromium-rich nickel-based alloys. As in cobalt alloys, during solidification, the HfC carbides precipitate in the interdendritic spaces of these nickel alloys, as script-like eutectic carbides mixed with matrix [6]. Thus, hafnium, the best known effect of which is the improvement of the behavior of superalloys in high temperature oxidation [7,8], in terms of oxidation rate or oxide scales adherence for instance, is also able to provide carbides potentially efficient for the mechanical reinforcement of superalloys.

Paradoxically, the hafnium carbides were rarely used to mechanically reinforce alloys for high temperature applications. HfC were however considered, thirty years ago, in cobalt-based hard metals [9] or in directionally solidified nickel-based alloys [10] for example. HfC were more used in molybden-based [11] or tungsten-based [START_REF] Chen | State-of-the-Art[END_REF] alloys. The goal was generally to improve the mechanical properties of these alloys. This was also the objective for the tensile properties [START_REF] Luo | [END_REF] or the creep resistance [14] of (molybden, tungsten)-based alloys, as well as of (tungsten, rhenium)-based alloys [START_REF] Park | Tungsten Tungsten Alloys[END_REF][START_REF] Park | [END_REF][17]. In contrast it seems that no significant results are available about the oxidation at high temperature of alloys containing great fractions of hafnium carbides.

Hafnium is a highly oxidable element. Its presence in high temperature alloys in quantities significantly higher than what is generally added to improve their hot oxidation behavior, may conversely threaten the alloy's resistance to oxidation at high temperature. This is notably true for uses at very high temperature at which the high thermal stability of HfC alloys can be better exploited. The purpose of the present work is to verify this point. The high temperature oxidation behavior of cast chromium-rich nickel-based alloys containing significant quantities of HfC carbides has been explored by performing oxidation tests at 1200°C for three alloys of this type.

EXPERIMENTAL

The work started by the elaboration of three Hf-rich nickel-based alloys containing 25 wt.%Cr, 0.25 or 0.50 wt.%C and 3.7 or 5.6wt.%Hf. Two supplementary ternary alloys with the same chromium and carbon contents were also synthesized to allow comparisons with the three former alloys in order to better analyze the effect of these high hafnium quantities on the hot oxidation behavior. All the alloys were elaborated by melting pure elements (Alfa Aesar, purity > 99.9wt.%) in a High Frequency induction furnace (CELES). Heating, melting, solidification and solid state cooling were realized under an inert atmosphere of 300mbars pure argon. Five compact-shaped ingots (weight: about 40 g) were obtained. They were cut for preparing metallographic samples for examining the as-cast microstructure, and parallelepiped samples for the thermogravimetry tests (about 9 × 9 × 3 mm 3 ). The metallographic samples were embedded in a cold resin mixture (resin CY230 + hardener HY956, ESCIL). They were polished with SiC-enriched papers with grade varying from 120 to 1200, cleaned by ultrasonic vibrations in water, and polished again with textile enriched with 1µm-alumina particles until a mirror-like state was obtained. The parallelepiped samples were polished with SiC-papers (120 to 1200-grit) and their edges and corners were smoothed using 1200-grit paper.

The thermogravimetry samples were oxidized, with mass gain recording, using a SETARAM TGA92 thermobalance. The whole thermal cycle was performed in dry industrial air (80%N 2 -20%O 2 ). It was composed of a heating at +20°C/min, an isothermal dwell of 46h at 1200°C and a cooling rate at -5°C/min. The mass gain files were first plotted versus time to specify the kinetic character (parabolic or not). The mass gain per surface unit area (m) was then plotted versus the square root of time for the classical determination of the parabolic constant k p . The mass gain files were also plotted and exploited according to the {m × dm/dt = k p -k v × m} method [18] to take into account the probable volatilization of chromia in order to avoid underestimating the parabolic constant k p . The chromia volatilization constant k v was determined by considering the slope of the linear part of the {(m × dm/dt) plotted versus (-m)} graph and the parabolic constant k p was determined by considering the ordinate at the origin of this same linear part.

Post-mortem characterization of the oxidized samples was done after oxidation test. They were first scanned using a simple office scanner to get macrographs illustrating the oxide spallation during cooling, and subjected to X-Ray diffraction (Philips X'Pert Pro diffractometer) for analyzing the natures of the external oxides remaining on surface. The oxidized samples were thereafter covered by an electrolytic nickel layer for protecting the external oxide from possible damage during cutting. They were cut in two parts, embedded in cold resin mixture and polished to obtain a mirror-like surface, as described above for the metallographic samples for as-cast microstructure observation.

All observations were done using a Scanning Electron Microscope (SEM) JEOL JSM-6010LA in Back Scattered Electrons (BSE) mode. The control of the general chemical compositions of the alloys was realized on the metallographic samples with the Energy Dispersion Spectrometry (EDS) apparatus equipping the SEM. The same device was used for identifying the carbides in the bulks of the as-cast samples and of the oxidized samples, the internal oxides in the subsurface, and the external oxide scales. This EDS device was also used to measure the chromium contents in extreme surface.

3

RESULTS

The Studied Alloys

The obtained chemical compositions of the studied alloys are given in Table 1 (measured by EDS). They are close to the targeted compositions. The determination by EDS was not possible for carbon because this element is too light and it is present in too low quantities in the studied alloys to be successfully analyzed by EDS. However, the obtained volume fractions of carbides showed that the targeted carbon contents were certainly well respected in each of the five alloys.

The as-cast microstructures of the alloys are illustrated by SEM/BSE micrographs in Figure 1. The two ternary alloys are composed of an austenitic dendritic matrix and of chromium carbides present in the interdendritic spaces. These carbides are of the M 7 C 3 stoichiometry. They are logically more numerous in the 0.50C-containing alloy than in the 0.25C-containing alloy.

The three Hf-rich alloys are also composed of a dendritic matrix and of interdendritic carbides. The matrix is a solid solution of nickel which contains the whole chromium, or the main part of the total chromium quantity if the alloy contains also chromium carbides. In contrast Hf seems to be absent in matrix, as shown by the EDS pinpoint measurements. The carbides are present, all in the interdendritic spaces, in higher quantity in the 0.50C-containing alloys than in the 0.25C-containing ones. These carbides are exclusively HfC (clearer than matrix in SEM/BSE mode) in the Ni-25Cr-0.25C+Hf and Ni-25Cr-0.50C++Hf alloys. This is in accordance with what was wished: the Hf and C weight contents were defined to have as many hafnium atoms as carbon atoms in these two alloys. In contrast, HfC are present together with chromium carbides in the Ni-25Cr-0.50C+Hf alloy which contains more carbon atoms than hafnium Hf atoms.

The Thermogravimetry Results

The obtained mass gain curves are given in Figure 2, on the left for the two alloys containing 0.25 wt.%C (the ternary alloy and the first Hf-rich one) and on the right for the three alloys containing 0.50wt.%C (the other ternary alloys and the two last Hf-rich alloy). All the curves are of the parabolic type, except in the case of the Ni-25Cr-0.50C++Hf, the mass gain kinetic of which seems becoming para-linear near the end of the 46 hours. The mass gain curves of the two ternary alloys are very close to one another. The presence of hafnium in the Ni-25Cr-0.25C+Hf alloy obviously led to a mass gain rate reduction by comparison with the corresponding ternary alloy. In contrast, the inverse effect is observed for the Ni-25Cr-0.50C+Hf alloy which oxidized slightly faster than the ternary Ni-25Cr-0.50C alloy. The oxidation rate of the Hf-richest Ni-25Cr-0.50C++Hf alloy is higher than for the Ni-25Cr-0.50C+Hf, which confirms this effect of hafnium for the 0.50C-containing alloys.

These mass gain files were all treated to specify first a linear k l value (slopes of the straight lines tangent to the mass gain curves at the beginning of the isothermal stage, Figure 2) and a k p value (classical method). The obtained values are given in the two first columns of results in Table 2. There is seemingly no systematic effect of the presence of hafnium on the value of k l since this one is lower for te 0.25C-containing Hf-rich alloy than for the corresponding Ni-25Cr-0.25C alloy while the inverse result is observed between the other alloys (0.50C without or with Hf). The same comments can be done about the classically determined k p values.

To take into account the mass loss by chromia volatilization, and thus to avoid underestimating the mass gain by oxidation and consequently the chromium-consumption, all curves were also treated according to the {m × dm/dt = k p -k v × m} method. In this type of representation (example of the Ni-25Cr-0.50C-3.72Hf in Figure 3), after an initial part characterized by very high values of the {m × dm/dt} term (on the right), a linear part (elongated cloud of points) appears in the graph for lower values of {-m}, as is to say for higher values of m. The corresponding values of real k p and of k v were red in the equation of the regression straight line drawn through the linear part of the graph (example in Figure 3). Before any comparison and interpretation these values of k p and k v need to be preliminarily verified. For this, they were tested. This started by drawing the theoretical curves calculated by using either the values of k p classically determined (mass gain plotted versus the square root of time), or the couples of values (k p , k v ) determined according to the {m × dm/dt = k p -k v × m} method. Thereafter the comparison of these two mathematical curves with the experimental curve plotted with the experimental mass gain values acquired for a given alloy, allows verifying whether the values obtained for k p (and of k v in the second case) well represent the real mass gain kinetic or not. Generally the curves plotted using the (k p , k v ) couples of constants' values better fitted the experimental curves than the curves plotted using only the classically determined k p constant. Indeed, the theoretic curve with k p and k v was generally almost superposed with the experimental one. Or at least, the end of the theoretic curve was parallel to the end of the experimental curve (example of the Ni-25Cr-0.50C-3.72Hf alloy in Figure 4).

The values obtained for k p and k v determined according to the {m × dm/dt = k p -k v × m} method are given in the two last columns of Table 2. The kp values are all globally twice the classical k p previously determined for the same alloys. They follow the same variations as the classical k p : decrease with Hf for the alloys containing 0.25 wt.%C and increase with Hf for the alloys containing 0.50 wt.%C. In contrast, the variation of k v with the presence of Hf is the same for the two groups of alloys: the volatilization constant increases from the ternary alloy to the Hfcontaining alloy and goes on increasing when the Hf content becomes higher.

Characterization of the Oxidized Surfaces

The oxidized samples were scanned on their two main faces. The macrographs of one of these two faces are presented in Figure 5. Obviously spallation occurred during cooling essentially for the two ternary alloys on which, after return to room temperature, the external oxide has been lost here and there. This is not the case for the three Hf-rich alloys for which the external oxide scale formed during the oxidation test remained over the whole surface. Hafnium had obviously here the same beneficial effect as usually observed for alloys containing more classical Hf contents.

The surfaces of the oxidized samples were also subjected to X-Ray Diffraction. Two examples are given in Figure 6 for one of the ternary alloys (Ni-25Cr-0.50C) and in Figure 7 for one of the Hf-rich alloys (Ni-25Cr-0.50C++Hf). All diffractograms showed that the main oxide present on the surface of the samples is chromia. However, for the Hf-rich alloys, HfO 2 oxides were also detected.

The samples were thereafter covered by electrolytic nickel, cut, embedded and polished. Despite the electrolytic nickel coating, cutting the oxidized samples provoked the loss of a part of the oxides initially present on the external surfaces. However, some oxides remained on surface. Their examination (SEM) and characterization (EDS) were thus possible to do on the metallographic cross-sections (examples in Figure 8). The ternary alloys are effectively covered by chromia. They classically present a carbide-free zone developed from the oxide/alloy interface. Chromia was also the main oxide present on the surface of the Hf-rich alloys but they also contained HfO 2 oxides. These later oxides are mixed with chromia in the external scale for some of them, or present in the subsurface as internal oxides. This confirms the results obtained by X-Ray Diffraction and given above. The oxidized Hf-rich alloys present another difference with the ternary alloys: no carbide-free zone developed from the alloy/oxide interface. In contrast with the chromium carbides of the ternary alloys, the HfC carbides did not disappeared in the subsurface; but some of them were locally oxidized. To finish one must also notice that more important internal oxidation seems having started in some locations in the subsurface of the Ni-25Cr-0.50C++Hf alloy.

Characterization of the Bulk Microstructures

The microstructures in the bulk of the Hf-rich alloys were also examined. They are illustrated by SEM/BSE micrographs in Figure 9, in which the initial as-cast microstructures are also reminded to facilitate comparisons. There was seemingly only a slight decrease in surface fraction of the HfC carbides in the three cases. In addition these carbides were slightly fragmented. Indeed the HfC carbides knew a limited loss of their initial script-like shape and only started to be replaced by aligned small round HfC carbides. However these two phenomena have progressed much slower than for TaC carbides in other alloys (cobalt-based for example) for the same conditions of temperature and duration.

DISCUSSION

Equiaxed nickel-based superalloys issued from classical foundry were seemingly never considered for the strengthening by hafnium carbides. An efficient reinforcement of such alloys may be though achieved by this way as this can be expected by looking to the microstructures obtained in this work. Indeed, interdendritic script-like HfC carbides mixed with matrix were successfully obtained. This was not initially evident since other script-like MC carbides, tantalum carbides TaC, did not precipitate in so high quantities in nickel-based alloys [4] although their carbon and tantalum contents were similar to the carbon and hafnium contents of the alloys studied here. It can be easily supposed that the mechanical strength at high temperature brought by other MC carbides in other austenitic equiaxed cast alloys (e.g. TaC in cobalt-based alloys [19]) will result from such microstructures for these HfCcontaining nickel-based alloys. But, before studying the mechanical behavior of these alloys at high temperature in hot air it was here verified whether the high Hf contents necessary to develop a dense interdendritic HfC carbides network (similar to the TaC network efficiently reinforcing some cobalt-based alloys) remains still compatible with sufficient oxidation resistance at high temperature.

The answer is that the presence of so high contents in hafnium is obviously not without consequences for these alloys. Indeed the kinetic was always parabolic but it was either accelerated or inversely slowed down, this depending on the carbon content. This may be explained by the fact that a higher carbon content may lead to a better exposure of the interdendritic carbides to oxidation (direct on surface and oxidation of Hf in the external scale, inwards diffusion of oxygen penetrating the alloys to form internal HfO 2 oxides), with consequently a slightly faster mass gain in the 0.50 wt.%C case than for the corresponding ternary alloys. When the carbon content is lower Hf may influence the oxidation rate as it usually acts when it is present in low quantities (decrease in oxidation rate). The presence of hafnium obviously influences chromia volatilization, an effect which is more difficult to interpret. But one can think in a first time that the presence of hafnium and its oxidation in HfO 2 may induce a destabilization of the growing chromia scale. This one may become more sensible to a new oxidation into CrO 3 , oxide which is gaseous at such very high temperature.

It is not surprising that the well-known effect of hafnium (when present in small quantities in superalloys) on the oxide spallation in case of cooling or varying temperature, was found for the alloys studied in this work. Hafnium obviously led to a significant improvement of the external oxide scale adherence, which can be of great interest in case of thermal cycling as it can be encountered in real industrial applications.

To summarize, the presence of 3.7 or 5.6 wt.% of hafnium in the HfC form and with an interdendritic repartition is not neutral for the high temperature oxidation behavior of these equiaxed cast chromia-forming nickel-based quaternary alloys. But the changes are not really significant for isothermal oxidation at 1200°C in air while there is in contrast a true improvement for the resistance against oxide spallation in conditions of cycling temperature. However it appeared necessary to keep an eye on the Hf-richest alloy which showed a tendency to an early start of linear oxidation, as revealed by the paralinear character of the mass gain near the end of the isothermal stage (after about 25 hours, Figure 2) as well as by the start of significant penetrating oxidation (Figure 8).

CONCLUSION

These Hf-rich nickel-based alloys have obviously kept the very good behavior in high temperature oxidation of the ternary alloys, thanks to the same high chromium content (25wt.%) but also to the not harmful effect to which the presence in so high quantity of the very reactive element hafnium possibly let think. The high temperature oxidation behavior was still chromia-forming, a character which seems to be not lost soon since the chromium contents, although decreased during these first 46 hours, were still high to ensure the perenity of this chromiaforming behavior. This, with the good volume fraction and morphological behavior of the reinforcing carbides, is favorable to a good sustainability of the high temperature properties of these new alloys. Indeed they are globally oxidation resistant but probably also corrosion resistant thanks to their high chromium contents, and creep resistant thanks to the good stability of the eutectic HfC carbides at high temperature, a creep-resistance which may be better than for TaC-reinforced cobalt-based alloys for which the fragmentation of the TaC was more severe for similar duration at the same temperature [20]. In contrast with similar cobalt-based alloys containing also high fractions of HfC carbides but not so resistant against high temperature oxidation [21], these alloys are thus ready to be tested at high temperature, without protective coating, to verify the later point: notably specifying their resistance against creep deformation.

FIGURE 1 .

 1 FIGURE 1. SEM/BSE micrographs of the as-cast microstructures of the five studied nickel-based alloys

FIGURE 2 . 11 - 2 FIGURE 3 .

 21123 FIGURE 2. The mass gain curves obtained; one the left: for the two 0.25C-containing alloys, and on the right: the three 0.50C-containing alloys

FIGURE 4 .FIGURE 5 .FIGURE 8 .FIGURE 9 .

 4589 FIGURE 4. Example of test of k p value classically determined and of a (k p , k v ) couple of values resulting from atreatment of the mass gain results according to the {m × dm/dt = k p -k v × m} method[18]; the best fit with the experimental curve is obtained with the (k p , k v ) curve (better parallelism of the ends of curves) (here: Ni-25Cr-0.50C+Hf oxidized at 1200°C)

TABLE 1 .

 1 Chemical compositions of the five nickel-based alloys (SEM/EDS measurements; +/-1wt.%)

	Elements	Cr	C	Hf
	Ni-25Cr-0.25C	25	supposed to be 0.25	/
	Ni-25Cr-0.25C	26	supposed to be 0.25	4.4
	+ HfC
	Ni-25Cr-0.50C	26	supposed to be 0.50	/
	Ni-25Cr-0.50C	26	supposed to be 0.50	4.8
	+ HfC
	Ni-25Cr-0.50C	25	supposed to be 0.50	6.6
	++ HfC

TABLE 2 .

 2 Values of the linear constant k l , of the parabolic constant k p and of the constant chromia volatilization constant k v , the two later ones being derived from the m × dm/dt = f(m) plots.

	Constants	k l (

× 10 -8 g/cm 2 /s) k p (× 10 -12 g 2 cm -4 s -1 ) k v (× 10 -10 g/cm 2 /s)

	method	slope at stage start	classic	m × dm/dt = f(-m)	m × dm/dt = f(-m)
	Ni-25Cr-0.25C	89	24	40	68
	Ni-25Cr-0.25C	37	12	36	126
	+ HfC
	Ni-25Cr-0.50C	25	15	32	93
	Ni-25Cr-0.50C	47	20	42	96
	+ HfC
	Ni-25Cr-0.50C	67	32	68	137
	++ HfC

ACKNOWLEDGMENTS

The authors wish to thank Pascal Villeger who has performed the X-Ray Diffraction runs.


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

Elodie Conrath
email: elodie.conrath@univ-lorraine.fr

High Temperature Oxidation of HfC-containing Chromiumrich Iron-based Alloys

Keywords: Iron-based alloys, HfC carbides, High temperature oxidation, Mas gain rates, Oxide scale characteristics, Internal oxidation

Three HfC-containing iron-based alloys were successfully elaborated by foundry. The alloys, all containing 25wt.%Cr, 0.25 or 0.50wt.%C and 3.7 to 5.6wt.%Hf, are characterized by a dendritic matrix and the presence of interdendritic script-like HfC carbides. These alloys were subjected to thermogravimetry test in air at 1200°C for 46 hours. The long exposure at this high temperature did not significantly modify the characteristics of the bulk HfC carbides. In contrast the presence of hafnium led to not parabolic oxidation kinetics and fast mass gain rates, by comparison with similar Hf-free ternary alloys. The scales formed all over the samples' surfaces were never composed exclusively of chromia and iron oxides or (iron, chromium) oxides were present in significant quantities. Some particles of hafnium oxides were also present in the external scales. A subsurface zone in which the HfC carbides were in situ replaced by hafnium oxides developed from the oxidation front. Catastrophic oxidation started in several locations and threatened to expand all around the samples beyond the stage duration. The high temperature oxidation behavior of these HfC-containing iron-based alloys appeared better than their cobalt-based homologues but worse than their nickel-based homologues, for the same conditions of oxidation test.

INTRODUCTION

Some applications at high temperature require refractory metallic alloys for their high ductility and toughness at all temperatures. Since such materials are intrinsically easy to oxidize -especially in hot conditions of servicethey must be enriched, either in bulk or in subsurface, by elements as aluminum, chromium or silicon [START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF]. When a good resistance against corrosion by molten substances is compulsory chromium is generally chosen since the developed chromia is often the most protective oxide layer in such conditions. Alloys rich enough in chromium for this purpose cannot contain significant quantities of aluminum and thus they cannot be strengthened by dense and fine precipitation of Ni 3 Al-type intermetallic particles such as the famous γ/γ' single crystal nickel-based superalloys [START_REF] Donachie | Superalloys: A Technical Guide[END_REF]. To resist mechanical stresses at high temperature chromium-rich alloys need to be constituted of coarse grains (achievable by foundry) and must contain particles reinforcing grain boundaries and interdendritic spaces, as well as particles or heavy atoms for the strengthening of the matrix itself. Carbides can be easily obtained in cast alloys by adding carbon and carbide-forming elements. Chromium already allows obtaining different types of chromium carbides, i.e. Cr 23 C 6 or Cr 7 C 3 which are the most common ones in alloys not too rich in carbon. They are for example the main reinforcing particles in cast chromium-rich nickel-based alloys used for glass-working [START_REF] Vasseur | Nickel alloy for centrifugal spinning apparatus for manufacture of glass fibers[END_REF]. Such carbides are unfortunately not stable enough at high temperature to insure resistance at elevated temperatures on long times. The TaC carbides show a better behavior in this field and they can be successfully obtained and exploited for the high temperature mechanical resistance in chromium-rich cobalt-based alloys. However, on the one hand such cobalt-based alloys are not very refractory since their fusion onset temperature is generally close to 1300°C and often lower, and on the other hand the initially script-like eutectic TaC carbides tend to get fragmented into aligned small and compact particles after less than 100 hours at very high temperature.

It was recently seen that other MC carbides -the HfC ones -appeared more resistant at high temperature against this fragmentation phenomena when obtained in chromium-rich cobalt-based alloys [4]. In parallel it was also observed in another earlier work [5] that MC carbides (the TaC ones) were successfully obtained with the script-like eutectic shape in chromium-rich iron-based alloys which appeared exceptionally refractory with fusion onset temperatures higher than 1400°C. It was thus envisaged to synthesize by foundry chromium-rich alloys containing quantities in carbon and hafnium high enough to obtain dense precipitation of HfC as main carbide phase.

This association of a very refractory Fe(Cr)-base and a strengthening HfC phase appears rather original. Indeed hafnium is generally considered for high temperature alloys as an element added with the typical amount of 1wt.%Hf or less for improving the behavior in high temperature oxidation [START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF]6] . This element was rarely used for high temperature mechanical strengthening purpose. One can cite few articles in which the formation of HfC carbides was promoted for the reinforcement by these carbides of cobalt-based alloys [7], directionally solidified nickel-based alloys [8], molybdenum-based alloys [9] or tungsten-based alloys [START_REF] Chen | State-of-the-Art[END_REF] alloys.

In the apparently single work dealing with the presence of HfC carbides in a chromium-rich iron-based alloy, it was found that script-like eutectic carbides are able to appear in the interdendritic spaces [START_REF] Berthod | [END_REF]. The purpose of this work is to synthesize again such alloys in order to test them in oxidation in air at a particularly high temperature (1200°C) for which such alloys can be destined according to their high refractoriness and to the advantageous morphologies and repartition of the very stable HfC carbides.

EXPERIMENTAL

Three Hf-containing iron-based alloys were synthesized by targeting the following compositions: 25 wt.%Cr, 0.25 or 0.50 wt.%C, and 3.7 or 5.6wt.%Hf. In parallel, to allow evidencing the influence of the previous high Hf contents on the high temperature oxidation behavior, two Hf-free ternary alloys containing the same chromium and carbon quantities were also elaborated. In each case pure elements (Alfa Aesar, purity > 99.9wt.%) were melted together in a High Frequency induction furnace (CELES). Heating, melting, solidification and solid state cooling took place inside a closed silica tube containing an inert atmosphere of pure argon (300mbars). Five compact 40gweighing ingots were obtained and, after return to room temperature, they were cut for obtaining samples of two types: metallographic samples devoted to the examination of the as-cast microstructures, and parallelepiped samples for the thermogravimetry tests (approximate dimensions: 9 × 9 × 3 mm 3 ). The samples for metallography examinations were embedded in a cold resin mixture (resin CY230 + hardener HY956, ESCIL) then polished. Polishing was done first by using SiC-enriched papers with grade varying from 120 to 1200. After cleaning using ultrasonic vibrations in water, the metallographic samples were finished using textile containing 1µm-alumina particles. A mirror-like state was thus obtained for all metallographic samples to allow their examination using optical and electron microscopes.

The parallelepiped samples destined to the oxidation tests in thermobalance were polished with SiC-papers (120 to 1200-grit). Their edges and corners were smoothed using 1200-grit paper. They were oxidized in a SETARAM TGA92 thermo-balance and their mass variations were recorded versus time. The applied thermal cycle, realized in dry synthetic air (80%N 2 -20%O 2 ), was composed of a heating at +20°C/min, of an isothermal stage of 46h at 1200°C and of a cooling rate at -5°C/min. The mass gain files were plotted versus time first to see if the kinetic is parabolic or not. The mass gain per surface unit area (m) was thereafter plotted versus the square root of time to classically determine the parabolic constant k p . At a so high temperature as 1200°C, when chromia is the major oxide to develop over the samples, its volatilization by re-oxidation into gaseous CrO 3 must be taken into account to do not minimize the oxidation rate when assessed by mass gain recording. The curves displaying a shape near parabolic and regular enough were then selected for attempting their analyze according to the {m × dm/dt = k p -k v × m} method [12] to specify the parabolic constant k p without underestimation and the chromia volatilization constant k v simultaneously. The curve obtained by plotting the m × dm/dt quantity versus -m (m being the mass gain per surface unit area) appears linear for values of times high enough (i.e. after the end of transient oxidation). The two kinetic constants can be then red in the equation of this straight line: chromia volatilization constant k v is the slope and the parabolic constant k p is the ordinate at the origin of this same linear part.

After the oxidation runs the thermogravimetry samples were characterized. This started by scanning the oxidized surfaces by using a simple office scanner (evaluating the possible oxide spallation occurred during cooling) and by performing X-Ray diffraction on the external oxide scales using a Philips X'Pert Pro diffractometer (identifying the natures of the external oxides present on surface). A thin gold layer was deposited all around the oxidized samples by cathodic evaporation, followed by an electrolytic layer of nickel thick enough in order to limit the loss of external oxide during cutting. After cutting in two parts, embedding in cold resin mixture and polishing to obtain a mirrorlike surface, the cross-sections were examined. This was done by optical microscopy (Olympus BX51 equipped with a digital camera driven by the Olympus software DP-Soft) and using a Field Emission Gun -Scanning Electron Microscope (FEG-SEM) JEOL J7600F (imaging in Back Scattered Electrons (BSE) mode and chemical analysis by Energy Dispersion Spectrometry (EDS)).

RESULTS

Chemical Compositions and Microstructure of the Studied Alloys

The final chemical compositions obtained for the elaborated alloys are given in Table 1 (EDS measurements). Globally the targeted chromium and hafnium contents are respected. Carbon being too light and present in too low quantity, it was not possible to measure its content directly. However it can be supposed that the carbon content was also well respected as one generally notes for similar alloys elaborated with the same procedure, this being confirmed in addition by seeing the carbides densities obtained.

The as-cast microstructures of the obtained alloys are illustrated by SEM/BSE micrographs in Figure 1. The matrix of the three Hf-containing alloys is dendritic. The carbides are either exclusively HfC carbides in the FeHf1alloy, consequence to its atomic ration between Hf and C close to 1. The two other Hf-containing alloys (HfC2-1 and HfC2-2), in which there is more carbon than hafnium in term of atomic contents, contain both HfC carbides (white in BSE mode) and chromium carbides (darker than matrix). When present the chromium carbides constitute an eutectic compound with the matrix. This is also the case for many of the HfC carbides which are scriptlike shaped and mixed with matrix in the interdendritic spaces. However there are also, in the two later alloys, several blocky polygonal HfC carbides, present here and there in the microstructure. These ones seemingly precipitated in the early stage of solidification (pre-eutectic carbides), simultaneously with matrix dendrites, and maybe even before the later ones.

Mass gain kinetics

The obtained mass gain curves are plotted in Figure 2 for the two 0.25wt.%C-containing alloys (Fe1 and FeHf1) and in Figure 3 for the three 0.50wt.%C-containing ones (Fe2, FeHf2-1 and FeHf2-2 alloy). The mass gain curve of the Hf-free Fe1 alloy is globally parabolic. The FeHf1 one has the same tendency for the first three quarters of stage duration (despite som small jumps due to probable local scale detachments), but it finishes by a more linear part. The Fe2 mass gain curve is not really parabolic but rather parabolic-linear. The mass gain curve obtained for the FeHf2-1 alloy was parabolic with a fast rate at the early beginning of the isothermal stage then it rapidly becomes parabolic-linear for the main part of the stage duration. Globally the mass gain rates of these four first alloys led to final mass gains of the same order of magnitude (about 3mg/cm² for the two Hf-free alloys and about 4mg/cm² for the two Hf-containing alloys). The FeHf2-2 alloy, in contrast, oxidized about 5 times faster, with a final mass gain of almost 20mg/cm². For this one the mass gain kinetic was first parabolic but with a high rate and it finished by a mass gain acceleration.

Considering the shapes of the obtained curves it appeared that only the Fe1 one was suitable for an exploitation of the mass gain files according to the {m × dm/dt = k p -k v × m} method. Its analysis, illustrated in Figure 4(A), led to values given in the first line of Table 2. These values, tested by plotting a mathematical curve plotted in Figure 4(B) ("kp-kv") for comparison with the experimental curve ("exp") with which it present a good parallelism for the two curve's ends. Thus the k p value issued from the {m × dm/dt versus -m}-plot is much better than the classical one which led to the third curve ("classic kp") which is very different from the experimental curve. The values of k p and k v -respectively 124 × 10 -12 g 2 cm -4 s -1 and 244 × 10 -10 g cm -2 s -1 -well correspond to what is generally obtained at such temperature when chromia cover the whole alloy surface (at least its main part) during oxidation [12,13]. The value obtained for kp, much higher than the classically determined one (47 × 10 -12 g 2 cm -4 s -1 ), clearly evidences the importance to take into account the volatilization of chromia, the rate of which is effectively not negligible).

Only the classical determination of k p was possible for the other Hf-free alloy (Fe2) and for the first part of the FeHf1 curve. The obtained values are close to the Fe1 one. The slope of the straight line drawn through the not parabolic mass gain curves obtained for the two 0.5wt.%-C Hf-containing alloys and plotted versus the square root of time, led to equivalent k p values very high, illustrating how fast these two alloys oxidized.

Surface characterization of the oxidized samples

The images obtained by scanning the oxidized sample before cutting and cross-sections preparation are displayed in Figure 5. This shows that oxide spallation during post-oxidation cooling were not severe since no really denuded zones can be seen. In contrast the irregular external face of the oxide scales let think that oxidation was not homogeneous all around the samples since thicker oxide islands obviously exist in some locations.

Concerning the natures of the oxides formed on surfaces it appears, thanks to the X-ray diffraction runs, that oxides of both Fe and Cr are present on the surfaces of all alloys: Fe 2 O 3 and Cr 2 O 3 well dissociated from one another, or (Fe,Cr) 2 O 3 mixed oxide, and in some cases FeCr 2 O 4 spinel. In the case of the Hf-containing alloys hafnium oxide (HfO 2 ) is additionally present, in small quantities. Two examples of such XRD diffractograms are given in Figure 6 for one of the Hf-free alloys and in Figure 7 for one of the Hf-containing alloys.

Characterization of the Oxides and Subsurfaces in Cross-Sections

The constitution of the external oxide scales in several oxides is already visible on optical micrographs (without preliminary etching) as illustrated in Figure 8. The observation using the optical microscope also allowed seeing that the external oxide scale can be, either thin and constituted of a single oxide or of the superposition of two continuous oxides, or thick and multi-constituted. The observations using the FEG-SEM brought additional information as the identification of the present oxides. One can se that spinel FeCr 2 O 4 and complex (Fe,Cr) 2 O 4 oxides are present under a continuous layer of chromia for the FeHf1 alloy (Figure 9), that Fe 2 O 3 and (Fe,Cr) 2 O 3 are on the contrary present over the chromia scale for the two other Hf-containing alloys (Figure 10 and Figure 11) and that more or less small particles of HfO 2 have formed in the external oxides involving Fe or Fe with Cr. The optical micrographs taken on the subsurfaces also indicated that internal oxidation may have occurred in all alloys but principally in the Hfcontaining ones. Some of the white particles seen in the subsurface (FeHf1 and FeHf2-1) alloys -and even in the core of the FeHf2-2 oxidized sample -are not HfC carbides but HfO 2 internal oxides.

The average thickness of the continuous oxides was measured (by avoiding the locations where very thick oxides locally formed) and is presented for all alloys as a graph in Figure 12. There is no very significant differences among the five results (by considering the standard deviation values) but there is a tendency to a thinner continuous scale for the Hf-containing alloys than for the Hf-free ones. The depth of internal oxidation (measured on optical micrographs for the reasons given above) is, in contrast, greater for the Hf-containing alloys than for the Hf-free ones (Figure 13). Furthermore the internal oxidation depth seems increasing with the Hf content in the alloy, this leading to the presence of dispersed hafnium oxides in the whole thickness of the Hf-richest alloy (FeHf2-2). Some EDS pinpoint measurements were performed in the subsurface close to the alloy-oxide interface. They showed that the chromium content had logically decreased from the initial values but remained still high (e.g. 19, 23 and 21wt.% in the FeHf1, FeHf2-1 and FeHf2-2 alloys respectively).

Microstructures Evolutions in the Bulk

The microstructures in the bulk of the oxidized Hf-containing alloys are displayed in Figure 14, Figure 15 and Figure 16 for the FeHf1, FeHf2-1 and FeHf2-2 alloys respectively, with SEM/BSE pictures taken for two magnifications (general microstructure and local details). Globally, the comparison with the as-cast microstructures displayed in Figure 1 shows that the microstructures have not significantly evolved, notably concerning the HfC carbides. These ones are still script-like shaped and their surface fractions seem to have remained the same. One can just remind the appearance of some hafnium oxides in the bulk of the Hf-richest alloy (for the two other internal oxidation occurred only in the first 250µm from the alloy/scale interface). In contrast, as for the Hf-free alloys, the chromium carbides, less stable at high temperature, have coarsened.

DISCUSSION

Thus, in chromium-rich iron-based alloys too, Hf additions in sufficiently great quantities allowed the development of numerous HfC carbides in the interdendritic spaces, mainly with a eutectic nature and a script-like shape but also with a blocky shape for some of them. This is a common point with the similar Hf-containing cobaltbased alloys [14] and Hf-containing nickel-based alloys [START_REF] Conrath | Oxidation of Metals[END_REF] which both presented also these repartition and shape of eutectic HfC carbides. However the presence of obviously blocky pre-eutectic HfC was met essentially in the cobalt alloys. As for these Hf-rich cobalt-based and nickel-based alloys previously elaborated by foundry and tested in oxidation at the same temperature the presence of so high quantities of hafnium influences the high temperature oxidation behavior of the corresponding Fe-25Cr-C base alloys. The kinetics were accelerated by the presence of hafnium, less than observed for the cobalt-based alloys [14] but more for the nickel-based alloys [START_REF] Conrath | Oxidation of Metals[END_REF]. By comparison with the nickel-based alloys for which the scale thicknesses were rather homogeneous, this is due to the formation of thick multi-constituted oxide islands. This is certainly the growth of the latter oxide islands which induced the linear mass gain contribution which added itself to the parabolic regime with as consequence a fast parabolic-linear kinetic.

Even if the high temperature oxidation behavior of the ternary Hf-free alloys was globally not excellent, it appeared that the presence of an interdendritic HfC network worsen this behavior: more iron oxide or (iron, chromium) oxides have formed and oxidation obviously became catastrophic in several locations. One can wonder whether the HfC network hindered the chromium diffusion towards the oxidation front, this effect being constant in time since HfC carbides obviously transformed on place into HfO 2 without liberating the interdendritic spaces close to the extreme surface. Indeed, they do not present a {carbide (HfC)} -free zone (or {oxide (HfO 2 )}-free zone) allowing the inwards development of a single-phased subsurface through which Cr atoms can diffuse in volume. The interdendritic HfC-carbides network, and thereafter the interdendritic HfO 2 -oxide networks, slowed Cr diffusion all along the oxidation test. This locally promoted the change of slow oxidation into fast catastrophic oxidation, phenomenon which may expand to the whole surface despite chromium contents remained rather high in subsurface (decreased only down to around 20wt.% against 25wt.% initially). This never occurred for the HfC-containing nickel-based alloys [START_REF] Conrath | Oxidation of Metals[END_REF], favored by the easy chromium diffusion through the nickel matrix. But, in contrast, this occurred sooner for the HfC-containing cobalt-based alloys [14] since these ones oxidized much faster all along the same isothermal stage, due to the globally less easy chromium diffusion through the cobalt-based matrix.

In contrast such carbide-free zone usually develops from the extreme surface in the case of TaC-containing alloys in the subsurface of which the tantalum carbides disappear over a depth increasing with time. In their case TaC dissociates into Ta and C which diffuse towards the oxidation front to be oxidized respectively into Ta 2 O 5 or CrTaO 4 close to the alloy/oxide interface for the first element and into CO x gases for the second one. For these TaCcontaining alloys this is probably the reason of their much better high temperature oxidation behavior insured by the chromium diffusion maintained easy by the absence of such obstruction.

Concerning the internal oxidation, this one can be attributed to the no diffusion of the immobility of the Hf atoms and to the easy diffusion of oxygen atoms through the ferritic matrix. No hafnium was found in the matrix and hafnium was exclusively present in carbides. Since the later are particularly stable no move from carbides to matrix is possible. During the oxidation test a part of oxygen diffused in the not compact ferritic matrix as well as along the interdendritic spaces, and formed preferentially hafnium carbides. HfC were thus replaced by HfO 2 in situ, over a depth increasing with time. The penetration of internal oxidation was obviously accelerated in case of dense HfC carbides as is to say for the FeHf2-2 alloy which presents simultaneously the highest carbon content and the highest hafnium content.

The HfC carbide network was obviously unchanged after almost fifty hours of exposure at 1200°C. This is due to the good stability at high temperature, which is much better than the chromium carbides one (chromium carbides have totally changed from their as-cast shape) and also better than the TaC carbides one (which generally tends to be more fragmented). Except in the subsurface over a depth of about 250µm no internal oxides replaced the hafnium carbides in the FeHf1 and FeHf2-1 alloys. In contrast, some HfO2 were seen among the HfC carbides in the bulk of the FeHf2-2 alloy. This may have (limited) effect on the mechanical behavior of the latter alloy, at high temperaturte as well as at medium and room temperatures (a little lack of ductility can be expected).

CONCLUSION

Such HfC-containing iron-based alloys are potentially very interesting since the HfC carbides present characteristics (interdendritic repartition, script-like shape mixed with the matrix… to add to the high refractoriness of such alloys) favoring high mechanical strength at high temperature such as creep-resistance. However, as the equivalent cobalt-based alloys with the same chromium, carbon and hafnium contents, they are obviously threatened by possible catastrophically oxidation, the first steps of which were observed here after only 46 hours of exposure to air at 1200°C. Modifications are required to improve their resistance against high temperature, such as Crenrichment of the bulk or surface Cr-enrichment by coatings (e.g. pack cementation). In a first time such improvements of the high temperature oxidation behavior will allow experimentally verifying their supposed high mechanical characteristics at high temperature. This can be the subject of further investigations.

FIGURE 2 .FIGURE 3 .FIGURE 4 .FIGURE 5 .FIGURE 8 .FIGURE 9 .FIGURE 12 .

 23458912 FIGURE 2. The mass gain curves obtained for the two 0.25C-containing alloys

FIGURE 13 .FIGURE 14 .FIGURE 15 .

 131415 FIGURE 13. Depth of the internal oxidation zone for all the alloys (average ± standard deviation)

TABLE 1 .

 1 Chemical compositions of the five nickel-based alloys (SEM/EDS measurements; +/-1wt.%)

	Elements	Cr	C	Hf
	Fe1	25.5	supposed to be 0.25	/
	FeHf1	25.7	supposed to be 0.25	3.9
	Fe2	24.5	supposed to be 0.50	/
	FeHf2-1	25.6	supposed to be 0.50	3.5
	FeHf2-2	27.2	supposed to be 0.50	4.9

TABLE 2 .

 2 Values of the linear constant k l , of the parabolic constant k p and of the constant chromia volatilization constant k v , the two later ones being derived from the m × dm/dt = f(m) plots. (*: values for allow qualitative comparisons, determined despite the curves are not really parabolic; **: the m × dm/dt = f(m) plots were not exploitable)

	Constants	k l (

× 10 -8 g/cm 2 /s) k p (× 10 -12 g 2 cm -4 s -1 ) k v (× 10 -10 g/cm 2 /s)

	method	slope at stage start	classic	m × dm/dt = f(-m)	m × dm/dt = f(-m)
	Fe1	114	34	124	244
	FeHf1	26	45*	/**	/**
	Fe2	34	47	/**	/**
	FeHf2-1	26	605*	/**	/**
	FeHf2-2	78	990*	/**	/**

FIGURE 1. SEM/BSE micrographs of the as-cast microstructures of the five studied iron-based alloys

ACKNOWLEDGMENTS

The authors wish to thank Sandrine Mathieu for the electron microscopy characterizations and Pascal Villeger for the X-Ray Diffraction runs.


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

Floriane Allègre
email: floriane.allegre6@etu.univ-lorraine.fr

Estelle Kretz
email: estelle.kretz@hotmail.fr

Influence of Titanium on the High Temperature Oxidation and

Keywords: Nickel-based alloys, Titanium, Chromium carbides, High temperature oxidation, Mass gain rates, Oxide scale characteristics, Internal oxidation

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INTRODUCTION

Nickel-based alloys represent a particularly important class of superalloys for high temperature applications [START_REF] Donachie | Superalloys: A Technical Guide[END_REF]. The most remarkable ones are the γ/γ' single crystals which offer very high levels of high temperature mechanical properties [START_REF] Gasko | [END_REF][3][4] and oxidation resistance [5][6][7][8]. Essentially devoted to constitute elongated pieces, these singlecrystal alloys are unfortunately not suitable to the compact geometries of many complex-shaped components for which other elaboration/shaping techniques must be preferred (conventional foundry, powder metallurgy…). However there are also polycrystalline nickel-based superalloys which can be employed for pieces with compact or complex shapes. In the case of these latter Ni-based alloys oxidation resistance at elevated temperature can be achieved with high amounts in chromium in their chemical compositions, instead aluminum is less efficient than Cr against corrosion by molten substances (CMAS, glasses…). Unfortunately, the protective external scale of chromia formed all around the Cr-rich superalloys tends to be re-oxidized into volatile species. This volatilization phenomenon is harmful for the sustainability of the oxidation resistance [START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF], especially when the working temperature is significantly higher than 1000°C.

Solutions for limiting chromia volatilization may be very interesting for keeping the longest time possible the chromia-forming behavior of superalloys. Adding reactive elements can be envisaged in this field, by hoping developing, together with chromia, another type of oxide whose presence may stabilize or protect Cr 2 O 3 from volatilization. Titanium, which is widely used for light alloys with high specific mechanical properties and oxidation resistance at moderate temperatures [START_REF] Pototzky | [END_REF][11][12][13][14], is one of such elements reacting easily with oxygen. Since this element seems to have never been evocated in association with the chromia volatilization phenomenon in the literature, it can be interesting to undertake a work to study whether Ti can have a positive effect in nickel-based chromia-forming superalloys by limiting the volatilization of Cr 2 O 3 at elevated temperature. This was done here at 1200°C, temperature high enough to promote significant volatilization of chromia, easy to characterize by thermogravimetry and to evidence a possible effect of titanium on this phenomenon.

EXPERIMENTAL

A series of five cast nickel-based alloys, containing first chromium in rather high quantity to favor a chromiaforming behavior and second interdendritic chromium carbides themselves favorable to an easy supplying of the oxidation front in Cr, were chosen to be sure to really obtain a chromia-forming behavior. Indeed it is known that Cr diffuses easily in the austenitic nickel matrix, much easier than in cobalt-based alloys for example. Chromium diffusion is furthermore helped by the presence of chromium carbides in the interdendritic spaces. High enough but still moderate levels of titanium were considered, in order to favor the observation of the possible effect of this element but without risking disturbing the microstructures of the model alloys. Three Ti-containing alloys were thus prepared by casting: Ni-25Cr-0.25C-1Ti ("NiTi1" alloy), Ni-25Cr-0.50C-1Ti ("NiTi2-1" alloy) and Ni-25Cr-0.50C-2Ti ("NiTi2-2" alloy). There were produced by melting in inert atmosphere (300 mbars of pure Ar) by using a CELES high frequency induction furnace. The melting of each mix of pure elements (Ni, Cr, Ti: from Alfa Aesar, purtity > 99.9%, and C: graphite, purity > 99.99%) was achieved in the water-cooled copper crucible of the furnace. After heating leading to total fusion, the liquid alloy was isothermally maintained during five minutes. Solidification and cooling down to ambient temperature were carried out in the copper crucible. For evidencing the possible effects of the presence of titanium two reference nickel-based Ti-free alloys with the same amounts in Cr and C were synthesized in parallel (Ni-25Cr-0.25C, "Ni1" alloy, and Ni-25Cr-0.50C, "Ni2" alloy), following the same elaboration procedure.

Five compact 40g-weighing ingots were thus obtained. They were cut for obtaining samples of two types: metallographic samples for the examination of the as-cast microstructures, and parallelepiped samples for the oxidation tests (approximate dimensions: 8mm × 8mm × 3 mm 3 ). The samples for metallography examinations were embedded in a cold resin mixture (resin CY230 + hardener HY956, ESCIL), and ground using SiC-enriched papers with grade varying from 120 to 1200. After cleaning using ultrasonic vibrations in water, they were polished using textile containing 1µm-alumina particles, to obtain a mirror-like state for all metallographic samples to allow their examination using electron microscopy.

The parallelepiped samples destined to the thermogravimetric tests were polished with SiC-papers (120 to 1200grit). Edges and corners were smoothed using 1200-grit paper. A SETARAM TGA92 thermo-balance was used for the oxidation runs for recording the mass variations versus time, in a 1.5L/hour flow of dry synthetic air (80%N 2 -20%O 2 ). The applied thermal cycle was composed of a heating at +20°C/min, an isothermal stage of 46h at 1200°C and of a cooling rate at -5°C/min to favor the keeping of the external oxide scales. The mass gain files were first plotted versus time to preliminarily see whether the kinetic is parabolic or not. The mass gain per surface unit area (m) was thereafter plotted versus the square root of time to determine the parabolic constant k p by applying an usual method. After having verified by X-ray diffraction that the external oxide scales were mainly chromia, its volatilization rate was assessed, as well as a more accurate kp constant (not minimized by the volatilization of chromia), by analyzing the mass gain files according to the {m × dm/dt = k p -k v × m} method [15], where m is the mass gain per surface unit area.

After the oxidation runs, the oxidized samples underwent metallographic characterization. This was done first by photographing the oxidized surfaces and measuring the fraction of surface where the external oxide scale obviously quitted the alloy, in order to characterize the amplitude of the oxide spallation when it occurred during cooling. Second, X-Ray Diffraction (XRD) runs were performed on the external oxide scales using a Philips X'Pert Pro diffractometer, for identifying the natures of the external oxides still present on surface.

A thin gold layer was thereafter deposited all around the oxidized samples by cathodic pulverization, followed by the electrolytic deposition of a nickel shell devoted to the protection of the external oxide during cutting. Cutting in two parts, embedding in cold resin mixture and grinding/polishing to obtain a mirror-like surface led to crosssectional samples ready to be examined by microscopy. Observations were done using electron microscopy (Scanning Electron Microscope JEOL JSM 6010LA), and Energy Dispersion Spectrometry (EDS apparatus attached to the SEM).

RESULTS

Chemical Compositions and Microstructure of the Studied Alloys

The chemical compositions of all the five alloys are given in Table 1. The wished contents in Cr and Ti were successfully reached, despite the high reactivity of Cr and especially Ti. However a few titanium oxides, nitrides and carbo-nitrides, identified by EDS spot analysis are present in the Ti-containing alloys, as illustrated in Figure 1 by SEM/BSE micrographs. In these ones one can see that the microstructures of the titanium-containing alloys are composed of an austenitic dendritic matrix of nickel-based solid solution (containing about 23 wt.%Cr for the 0.25C-containing alloy, and around 21.5 wt.%Cr for the two 0.50C-containing alloys) in which almost all the titanium atoms are present in solid solution, and of interdendritic chromium carbides which were identified by spot EDS analysis as being Cr 7 C 3 .

Mass gain kinetics

The obtained mass gain curves are plotted in Figure 2 for the two 0.25wt.%C-containing alloys (Ni1 and NiTi1) and in Figure 3 for the three 0.50wt.%C-containing ones (Ni2, NiTi2-1 and NiTi2-2 alloys). All the mass gain curves are globally parabolic and the determination of the classical k p values did not pose any problem. Since the external oxide scales were essentially composed of chromia, as this will be evidenced further, the mass gain files were also plot as m × dm/dt versus -m in order to determine the values of the two constants describing the oxidation kinetic. This was remarkably successful since the {m × dm/dt versus -m} plots led to linear parts (top of Figure 4 for the NiTi1 alloy, Figure 5 for NiTi2-1 and Figure 6 for NiTi2-2), the equations of which allowed determining the values of the volatilization constant (slope of the regression straight line) and the real parabolic constant (ordinate at the origin). The linear k v constant characterizing the rate of volatilization of chromia, and the real parabolic k p constant (not underestimated since the mass loss in chromium by leaving CrO 3 is then taken into account) were then determined (Table 2). As illustrated by the graphs presented in the bottom parts of the Figures 4, 5 and6, these values were validated by the very good correspondence between the theoretic mass gain curves calculated from these values of k p and k v (green curves) and the experimental mass gain curves (black curves). The theoretic curves plotted with the conventionally determined k p (red curves) do not match so exactly. Additionally, the linear constant k l associated to the transient oxidation at the beginning of the isothermal stage was also determined (on the classical m= f(t) plot).

The three Ti-containing alloys oxidized faster than the two Ti-free alloys. This is graphically visible and confirmed by the values of the parabolic constants. Furthermore, the volatilization constants determined for the Ticontaining alloys seem to be globally higher than for the Ti-free alloys (average k v of 95 × 10 -10 g cm -2 s -1 for the Ti-containing alloys against 81 × 10 -10 g cm -2 s -1 for the Ti-free ones).

Oxide spallation during cooling

After return at room temperature, it is obvious that some of the oxidized Ti-containing samples have lost only small parts of their external oxide scales, by comparison with the two ternary Ti-free alloys (Figure 7). Since the Ticontaining alloys did not lose anywhere their external oxide scale on the whole thickness the metal surface was visible and image analysis led to 0% of surface having lost the external oxide scale, even if it is obvious if some small parts of oxide were lost. This is different for the Ti-free alloys for which areas having lost their external oxide were clearly visible: they represent (average values for both main faces) 29% and 17% for the Ni1 and Ni2 alloys respectively.

The start of spallation during the cooling phase can be easily observed on the {mass versus temperature}-curves. Figure 8 (the two Ti-containing and Ti-free 0.25C-containing alloys) and Figure 9 (the two Ti-containing and Ti-free 0.50C-containing alloys) show that spallation started at 688°C for the NiTi1 alloy (i.e. a little later than the Ni1 one: 832°C), at 702°C and 737°C for the NiTi2-1 and NiTi2-2 alloys respectively (later than the Ni2 alloy: 817°C). In addition the mass variations along the whole thermal cycle were much more limited by comparison with the two ternary alloys (for which the total mass variation is negative!).

Analysis of the external oxides and of the subsurfaces

X-ray diffraction was carried out on the oxidized surfaces of the five alloys (example of NiTi1 in Figure 10). Unsurprisingly, chromia is present over all the oxidized surfaces but there is also a significant presence of TiO 2 . The two oxides are shown in the cross-sectional SEM/BSE micrographs given in Figure 11 for NiTi1, and in Figure 12 for NiTi2-1 and NiTi2-2, and the results of the measurements of the oxide scale thickness or carbide-free zone depth are given in Table 3 (average and standard deviation of five values). It appears that all alloys are covered by a continuous chromia layer (identified by EDS) which seems thicker for the Ti-containing alloys than for the Ti-free ones. Indeed the chromia scales formed over the Ni1 and Ni2 alloys (around 22 µm) are 10µm lower than the ones formed over the NiTi1, NiTi2-1 and NiTi2-2 alloys (around 33µm). The chromia scale was rather porous in all cases but maybe more for the Ti-containing alloys than for the Ti-free ones. In addition, for the Ti-containing alloys, a seemingly continuous outer layer made of TiO 2 exists and its thickness tends to increase with the Ti content in alloy (around 3µm for NiTi1 and NiTi2-1 against 5µm for NiTi2-2). This outer TiO 2 oxide scale seems having been lost in some locations. This can be due to metallographic preparation but it is also possible that parts of TiO 2 were also lost during the cooling, spallation detected on the cooling part of the {m versus temperature}-plotted thermogravimetry curves Figure 8 and Figure 9) and visible on the oxidized surfaces of the samples (Figure 7). Under the external scales carbides seemingly disappeared from the oxidation front, over a depth which is lower for the 0.50C-containing alloys than for the 0.25C-containing ones. These carbide-free zones tend to be a little deeper for the Ti-free alloys than for the Ti-containing alloys for a given carbon content. As expected they are impoverished in Cr and Ti (the ternary alloys only in Cr), as illustrated by the concentration profiles given in Figure 13 (0.25Ccontaining alloys) and Figure 14 (0.50C-containing alloys). The titanium contents in the matrix are equal to 0 in the extreme surface and they increase progressively when one goes deeper, this evidencing the diffusion of Ti towards the oxidation front. The same classical observation can be done for chromium. Globally the Cr-depleted depths correspond to the carbide-free zones while the titanium-depleted depths are more extended inwards. One can notice that the Cr content in extreme surface tends to be lower for the Ti-containing alloys than for the Ti-free ones, demonstrating that the former ones seem being threatened by the loss of the chromia-forming behavior than the latter ones.

DISCUSSION

Despite its high reactivity with oxygen and nitrogen titanium was thus successfully introduced in these base ternary alloys during conventional casting. It is true that even pure argon may contain extremely low quantities of O or N but high enough to allow Ti reacting with them. A few titanium oxides, nitrides or carbo-nitrides were present but they did not perturb the study. Titanium was almost exclusively present in solid solution without forming carbides, showing that TiC are less stable than chromium carbides in these nickel-based 25wt.%Cr-containing alloys.

The isothermal oxidation kinetics of the three Ti-containing alloys were of a parabolic type without any jumps, as the corresponding Ti-free ternary alloys. By comparison by these latter ones the mass gains of the Ti-containing alloys were significantly faster. This can be due to the participation of Ti in the formation of oxides (TiO 2 both internal and external), but it is also true that the external chromia scales were thicker than for the ternary alloys. Furthermore chromia seemed more porous than the ones formed over the ternary alloys. The reasons of these two differences are not evident and needs to be deepened. TiO 2 oxides formed internally and also externally. This second part of oxides is located outside and constitutes a thin layer which seems isolating chromia from hot air. Unfortunately this TiO 2 external layer covering the chromia scales obviously did not protect it from contact with the oxidizing atmosphere since the chromia volatilization constants of the Ti-containing alloys tend to be not lower but higher than for the Ti-free alloys. It is true that the outer TiO 2 scale obviously contains defects (pores…) which do not allow it totally isolating chromia from outside. The chromia scales observed on the Ti-containing surfaces are thicker than the ones formed over the Ti-free alloys, not because lower volatilization rates but because faster oxidation. This is confirmed by the consumption of chromium coming from the alloys which is greater for the Ti-containing alloys than for the Ti-free ones, as proven by the Cr profiles presented in Figure 13 and Figure 14. Furthermore, the average Cr gradients in subsurface are higher for the Ti-containing alloys than for the Ti-free ones, this suggesting that, even if the final mass gain rates were higher in the first case than in the second case, Cr diffused easier in absence of Ti in the matrix. Such an obstruction effect of the Cr diffusion due to Ti may threaten the chromia-forming behavior.

Thus, the presence of 1 or 2 wt.%Ti in the model ternary alloys was rather detrimental for the high temperature oxidation behavior of the alloys, but benefit was observed concerning the maintenance of the external scales over the samples during the cooling: spallation started later (i.e. at a lower temperature) during the post-isothermal stage cooling and the final mass loss was much lower than for the Ti-free alloys. It is not really possible to say presently if this due to the TiO 2 outer oxide scale or to a possible pegging effect (e. g. very small particles involving Ti favoring the adherence of chromia on the alloy). This positive influence of Ti can be of great importance in case of thermal cycling.

CONCLUSION

The presence of titanium in chromium-rich nickel-based alloys is not neutral for the oxidation behavior at high temperature. Here studied in the case of alloys based on ternary alloys of the {Ni-25Cr-0.25 or 0.50}-type, known to be chromia-forming and well resistant against oxidation in air even at very high temperatures, one clearly saw differences in terms of mass gain rates, oxide thickness, subsurface Cr-depletion, resistance against spallation and even of carbide-free depths. Globally Ti behaved here as rather detrimental in isothermal conditions and, in contrast, as rather beneficial in varying temperature conditions. In contrast, with what was observed for the Ti-free alloys (e.g. slower mass gain for 0.5wt.%C than for 0.25wt.%C) it was also pointed out that the oxidation behavior of the Ti-containing alloys did not globally depend on the carbon content (except of course the carbide-free depth, as usually observed for carbides-containing cast alloys).

The reasons of these effects of titanium, to confirm for other temperatures and durations of oxidation and other cooling rates, will be deeper studied in future investigations.

2 FIGURE 1 .FIGURE 2 .FIGURE 3 .FIGURE 4 .FIGURE 5 .FIGURE 6 .

 2123456 FIGURE 1. SEM/BSE micrographs of the as-cast microstructures of the three Ti-containing Ni-based alloys

FIGURE 7 .FIGURE 8 .FIGURE 13 .FIGURE 14 .

 781314 FIGURE 7. Aspects of the oxidized alloys after return to room temperature

TABLE 1 .

 1 Chemical compositions of the five nickel-based alloys (SEM/EDS measurements; +/-1wt.%)

	Elements	Cr	C	Ti
	Ni1	25.1 ±0.2	nominally 0.25	/
	NiTi1	25.4 ±0.1	nominally 0.25	1.0 ±0.1
	Ni2	25.2 ±0.2	nominally 0.50	/
	NiTi2-1	25.7 ±0.1	nominally 0.50	1.1 ±0.1
	NiTi2-2	25.5 ±0.3	nominally 0.50	2.1 ±0.1

TABLE 2 .

 2 Values of the linear constant k l , of the parabolic constant k p and of the constant chromia volatilization constant k v , the two later ones being derived from the m × dm/dt = f(m) plots.

	Constants	k l (

× 10 -8 g/cm 2 /s) k p (× 10 -12 g 2 cm -4 s -1 ) k v (× 10 -10 g/cm 2 /s)

	method	slope at stage start	classic	m × dm/dt = f(-m)	m × dm/dt = f(-m)
	Ni1	89	24	40	68
	NiTi1	34	51	86	110
	Ni2	25	15	32	93
	NiTi2-1	21	54	76	73
	NiTi2-2	24	51	85	101

TABLE 3 .

 3 Values of the linear constant k l , of the parabolic constant k p and of the constant chromia volatilization constant k v , the two later ones being derived from the m × dm/dt = f(m) plots.

	Thicknesses or depths (µm)	TiO 2	Cr 2 O 3	Carbide-free zone
	Ni1	/	19.6 ±2.1	181 ±11
	NiTi1	3.5 ±0.3	32.8 ±0.5	156 ±8
	Ni2	/	23.7 ±2.4	124 ±9
	NiTi2-1	2.3 ±0.3	32.9 ±1.8	123 ±6
	NiTi2-2	4.6 ±0.7	33.1 ±1.4	114 ±3

ACKNOWLEDGMENTS

The authors wish to thank Mélissa Ritouet and Thierry Schweitzer for their technical help.