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Oxidation behaviour of a Ti–46Al–1Mo–0.2Si alloy: the effect of Mo addition and alloy microstructure
Materials Science and Engineering A284 (2000) 138 – 147 www.elsevier.com/locate/msea
Oxidation behaviour of a Ti–46Al–1Mo–0.2Si alloy: the effect of Mo addition and alloy microstructure P. Pe´rez a,b,*, J.A. Jime´nez b, G. Frommeyer c, P. Adeva b b
a Institute for Health and Consumer Protection. JRC of the European Commision, I-21020 (Va), Ispra, Italy Centro Nacional de In6estigaciones Metalu´rgicas, CENIM, (CSIC). A6da. Gregorio del Amo 8, 28040, Madrid, Spain c Max-Planck-Institut fu¨r Eisenforschung GmbH, Max-Planck-Strabe 1, D-40237 Du¨sseldorf, Germany
Received 14 April 1999; received in revised form 21 January 2000
Abstract The influence of molybdenum addition and alloy microstructure on the oxidation behaviour of a Ti – 46.8Al – 1Mo –0.2Si alloy was studied in air at temperatures ranging from 600 to 900°C. The alloy produced by arc melting exhibited a structure of coarse lamellar grains in as-cast condition that transformed to a duplex microstructure after hot extrusion at 1300°C. Oxidation rate and scale spallation resistance were not affected by the type of microstructure. The effect of molybdenum addition was related to the formation and thickening of a protective continuous nitride layer. The low mass gain measured can be associated with the low oxidation rate of the nitride layer. Finally, the poor spalling resistance of the oxide scale at 900°C was attributed to the mismatch in thermal expansion coefficients of the oxide scale and the nitride layer. © 2000 Elsevier Science S.A. All rights reserved. Keywords: g-TiAl based intermetallic alloy; High temperature oxidation; Microstructure; Alloying
1. Introduction The research of structural materials for high-temperature applications has led to the consideration of gTiAl base alloys for aviation gas turbines or automotive engine components, since this intermetallic compound combines low density with high specific strength and stiffness up to 800°C. However, low ductility at room temperature, poor fracture toughness and insufficient oxidation resistance at temperatures above 700°C have limited the use of g-TiAl based alloys. Attempts for improving these properties have been made through additions of specific alloying elements and microstructural modifications. Several investigations have shown that V, Mn, and Cr generally increase the ductility [1], and elements like Mo, Nb, Hf, Zr, Cl or W improve the oxidation behaviour [2 – 8]. However, each of these alloying elements improves a specific property but can be detrimental to the other properties. Therefore, it is difficult to balance good mechanical properties and oxidation resistance only by alloying elements. * Corresponding author.
Changes in the microstructural morphology have been also reported to affect strongly the mechanical properties of g-based alloys [1]. For two-phase g-TiAl alloys, different microstructures can be generated depending on processing route and/or subsequent thermal treatments, i.e. fully lamellar, nearly lamellar, duplex or near g [1]. Generally, the duplex structure is desirable for ductility, but poor for toughness and creep resistance. The opposite is true for the fully lamellar structure. At this point, an adequate ternary element addition and control of alloy microstructure could be appropriated to improve both, mechanical properties and oxidation resistance of these alloys. However, the effect of type of microstructure on the oxidation behaviour has not been thoroughly investigated, as indicated by the controversy existing in the literature about this point [9–13]. Gil et al. [9,10] reported for a two-phase g/a2TiAl alloy oxidized at 900°C in air, that the oxidation rate and type of the formed scale depend on the type of microstructure and distribution of both phases. Cast material with a fine lamellar microstructure formed a thin protective scale instead of the thicker rutile-based scale established on heat treated material, whose mi-
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crostructure consisted of large grains of g-TiAl and coarse a2-Ti3Al. Recently, Pe´rez et al. [14] found important differences in the oxidation behaviour of fullylamellar and duplex Ti – 46Al – 1Cr – 0.2Si alloy. At temperatures ranging from 700 to 900°C, the duplex microstructure presented a higher oxidation rate than the fully-lamellar. On the other hand, Shemet et al. [15], for TiAl oxidized in argon/oxygen at 900°C and Haanappel et al. [12,13] for Ti – 48Al – 2Cr oxidized in air at 700 and 800°C, reported that oxidation rate was unaffected by the alloy microstructure. The aim of this study is to determine the influence of the type of microstructure, i.e. fully lamellar in as-cast condition and duplex after extrusion, on the oxidation behaviour of TiAl based alloys containing small additions of Mo. This work is complementary to a previous one on the influence of the microstructure on the oxidation behaviour of alloys containing small amounts of Cr (1 at.%), which is detrimental for the oxidation resistance [14]. In that case, it was found a considerable effect of the alloy microstructure on the oxidation behaviour. In the present work, the oxidation behaviour of fully-lamellar and duplex Ti – 46.8Al–1Mo– 0.2Si alloy was studied. Oxidation kinetics have been evaluated in air at temperatures ranging from 600 to 900°C.
2. Experimental A Mo-alloyed two-phase g/a2-TiAl alloy with composition (in at.%) Ti – 46.8Al – 1Mo – 0.19Si was produced in 50 kg ingots by arc skull melting. As-cast ingots exhibited a coarse-grained microstructure. This microstructure was refined by hot extrusion of billets 75 mm in diameter and 120 mm in length in the a + g field at 1300°C. The reduction in area of the extruded rods was about 7:1. Microstructure present in the as-cast material and after extrusion was characterized using X-ray diffraction, optical and scanning electron microscopy (SEM). The phases present and the volume fraction of each phase were determined by X-ray diffraction. The diffraction patterns showed the presence of reflections of both phases g-TiAl and a2-Ti3Al. The volume fraction of each phase was calculated from the integrated intensities of (201), (022), (220), and (203) peaks of the a2-Ti3Al, and (002), (200), (202), and (220) peaks of g-TiAl. A decrease in the amount of the a2-Ti3Al from 20 vol.% in the as-cast material to 10 vol.% after hot extrusion was found. This value agrees with that predicted by the equilibrium diagram, about 10 vol.% [16]. Oxidation tests were carried out on rectangular coupons of 10× 10×1 mm3 cut from the billets by electrospark erosion. The surfaces were abraded on successively finer silicon carbide papers, then mechani-
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cally polished with 1 mm diamond paste and finally ultrasonically cleaned with acetone. Oxidation kinetics in the temperature range from 600 to 800°C were determined by discontinuous thermogravimetry. The actual test temperature was verified by a thermocouple placed close to the sample. After exposure, samples were taken out of the furnace and left to cool down, then weighed using a balance with a resolution of 910 mg. Characterization of oxidation products was performed on samples isothermally oxidized for short and intermediate exposure times, as well as on samples used for the oxidation kinetics determination. Cross-sections of samples were prepared by conventional metallographic technique. To prevent scale loss during metallographic preparation sample surfaces were successively coated with a thin gold layer (by sputtering) and a thicker layer of copper (electrolytically deposited). Cross-sections were prepared by conventional metallographical techniques. Surfaces and cross-sections of the oxidized specimens were studied by SEM. Phase identification of the oxide scale was performed by X-ray diffraction (XRD), energy dispersive X-ray microanalysis (EDX) and wavelength dispersion X-ray microanalysis (WDX).
3. Results
3.1. Microstructure As-cast alloy presented a microstructure consisting of very coarse lamellar grains with a size of several hundred micrometers, as shown in Fig. 1(a). This lamellar structure was partially transformed by extrusion at 1300°C into a duplex microstructure consisting of single-phase g grains, ranging from 5 to 20 mm in size, and finer grains, as shown in Fig. 1(b). Observations in the extrusion direction showed that both single-phase g and two-phase g/a2 grains were grouped in alternating bands ranging from 60 to 150 mm in width (Fig. 2). The origin of such microstructure has been attributed to the segregation of aluminium during the peritectic solidification [17,18]. However, in the present work the development of this stripped microstructure must be related to the precipitation of small amounts of a Mo-rich phase, as confirmed by EDX microanalysis. These Morich particles, probably Ti2AlMo as found by XRD (see Fig. 3), are heterogeneously distributed on the microstructure, as observed in Fig. 4(a and b) corresponding to bands of coarse and fine grain, respectively. Mo-rich precipitates were only present in fine grain regions (see Fig. 4(b)). This non-homogeneous distribution results from segregation of molybdenum during the casting process.
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As Mo-rich particles were not found in the as-cast material, it was concluded that they precipitated before the extrusion, during heating. In the extrusion process, these particles pinned the grain boundaries of the recrystallized grains, preventing the grain growth during and after extrusion. However, as distribution of these
Fig. 3. XRD spectra of non-oxidized extruded alloy.
Fig. 1. Optical micrographs showing the microstructure of (a) as-cast material and (b) extruded material.
Fig. 4. Micrographs of the microstructure in coarse (a) and fine grain regions (b) of the extruded alloy. It is noticed the presence of Mo-rich precipitates exclusively at grain boundaries of fine grain regions.
Fig. 2. SEM micrograph of the extruded alloy showing the microstructure in the extrusion direction.
particles is heterogeneous, in those regions with a low volume fraction of particles the grain growth will not be inhibited, resulting in the stripped microstructure.
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3.2. Oxidation kinetics Oxidation kinetics of the Ti – 46.8Al – 1Mo –0.2Si alloy have been evaluated in air at temperatures ranging from 600 to 900°C by discontinuous thermogravimetry. The mass gain– time curves for both extruded and as-cast material are represented in Fig. 5. The general oxidation behaviour of the two microstructures, i.e. duplex and lamellar, is quite similar in the whole temperature range tested. At 600°C, the mass gains were negligible and the curves for both materials can be superimposed. At 700 and 800°C, the oxidation rate of both as-cast and extruded alloy tended to decrease with progressing oxidation until a steady state was reached. At 900°C, the kinetics were affected by an extensive scale spallation already after short exposure times, 1 and 10 h for as-cast and extruded material, respectively.
Fig. 6. Surface morphology of the scale formed after 160 h of oxidation at 600°C, showing microstructure of (a) as-cast and (b) extruded alloy (SEM micrographs).
3.3. Surface morphology
Fig. 5. Mass gain-time curves for both as-cast and extruded alloy oxidized in air at: (a) 600 and 700°C. (b) 800 and 900°C.
The thin scale formed after testing at 600°C revealed the microstructure of the alloy because dark and bright zones were formed above the g-TiAl and a2-Ti3Al phases, respectively (see Fig. 6). EDX analysis evidenced that the dark zones consisted of an Al-rich oxide, likely alumina, while the bright zones consisted of an Al-rich oxide covered by a very porous and discontinuous Ti-rich oxide, probably rutile. At 700°C, an alumina-rich layer formed at the surface from the early stages of oxidation. For longer exposure times, the rutile content of the scale increases and small well-developed rutile crystals commenced to appear progressively on the surface, as shown in Fig. 7. No differences between as-cast and extruded material were observed. At 800 and 900°C, the morphology of the scale was completely different from that described before. The outer layer consisted of the rutile phase, even in the earlier stages of oxidation. The size of the rutile crystals increases with increasing temperature and exposure time, as shown in Fig. 8.
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3.4. Cross-sectional obser6ations
Fig. 7. SEM micrographs showing surface scale morphology after 150 h at 700°C; (a) as-cast alloy, (b) extruded alloy.
Cross sectional examination revealed only small differences between the scale formed on the as-cast and extruded material. A thin uniform scale developed on both materials in the whole temperature range studied, as shown in Fig. 9. At 600°C, this scale was too thin to be analyzed. At 700°C, the scale formed after 50 h of exposure was very thin (B 1 mm). It consisted of two layers, an outer alumina-rich layer and an inner rutilerich layer. Low scale growth rate was observed during further oxidation. Thus, after 150 h of exposure an oxide scale 3 mm thick was observed, in which the inner part has a thickness of 2 mm and the outer part a thickness of 1 mm (see Fig. 10). A nearly continuous white Ti-rich phase appeared at the scale/alloy interface. At 800 and 900°C, the scale pattern for short oxidation times was practically identical to that observed at 700°C for longer times. The major difference concerned the development of a thin outer rutile layer (Fig. 11). The Ti-rich phase at the scale/alloy interface was identified by XRD as titanium nitrides of the type Ti2AlN (Fig. 12). In addition, also very small peaks of TiN (not observed in Fig. 12) were detected for both microstructures. After 160 h at 800°C, the scale was 10 mm thick and all the layers composing the scale thickened (Fig. 13). The intermediate alumina-rich layer was 2 mm thick and presented an apparently dense microstructure. Nevertheless, this alumina layer was not sufficiently protective because it prevented neither the
Fig. 8. SEM micrographs of the oxide scale morphology developed on as-cast (a, b) and extruded alloy (c, d) after 160 h of exposure at 800°C (a, c) and 900°C (b, d).
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exposure times at 900°C (Fig. 15). EDX analysis suggested that nitride layer was composed mainly of Ti2AlN (see Table 1). In addition, it was checked that sub-scale region contained a small amount of oxygen, about 3%.
4. Discussion The results obtained in the present work show that the mechanism controlling the oxidation of the Ti– 46Al–1Mo–0.2 Si alloy was very similar for both microstructures, i.e. fully-lamellar and duplex. The oxidation rate at 600, 700, 800 and 900°C was practically identical for both microstructures. At 600°C, the low oxidation rate suggests the formation of a continuous healing alumina-rich film over the entire surface. Above the alumina-rich layer situated at a2-Ti3Al lamellae appeared a porous rutile layer, probably formed during the initial stages of oxidation. The
Fig. 9. Cross-sections of the scales formed on (a, c, e) as-cast and (b, d, f) extruded alloy after 150 h 700°C (a, b), 160 h 800°C (c, d) and 1 h 900 (e, f) (SEM micrographs).
Fig. 11. Cross-section of the scale formed after 10 h at 800°C on the extruded alloy (SEM micrograph).
Fig. 10. Cross-sections of the scale formed after 150 h at 700°C on the extruded alloy (SEM micrograph).
outward diffusion of titanium nor the inward diffusion of oxygen. This resulted in a slow but continuous growth of the outermost and inner rutile-rich layers. The coalescence of the innermost titanium nitrides led to the formation of a continuous layer, 1 mm in thickness, at the scale/alloy interface. A SEM micrograph of the sample presented in Fig. 13 after etching is shown in Fig. 14. This figure revealed the formation of an Al-rich a2-Ti3Al free zone beneath the nitride layer. Both, the Al-rich a2-Ti3Al free zone and the nitride layer significantly thickened after longer
Fig. 12. XRD pattern of oxide scale formed on the extruded alloy after 24 h of oxidation at 800°C.
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At 700°C, the low oxidation rate measured in the early stages of oxidation for both alloys suggests that an alumina-rich film formed in this period. As oxidation proceeds, surface morphology observations showed the presence of rutile crystals, indicative of outward titanium transport through the scale. The appearance of these rutile crystals seems to account for the increase of the oxidation rate measured in the mass gain curves. At this stage, the alumina-rich scale left to be protective as proves the development of an inner thicker rutilerich layer. The formation of this layer implies inward oxygen through the outer alumina-rich layer. The decrease of the oxidation rate for longer times suggests that a more protective layer started to control the oxidation. It could be possible that thickening of the intermediate alumina-rich layer with time leads to a more compact and protective scale. Nevertheless, one might assume also that a new more protective layer starts to control the oxidation process. This new layer could correspond with the nearly continuous titanium nitride layer detected at the scale/alloy interface after longer times. The formation of a continuous nitride layer and further thickening should occur for long term exposures as it has been described for the oxidation of Ti–48Al–2Cr–2Nb alloy in air at 700°C [25]. Fig. 13. Cross-sections of the scales formed after 160 h of exposure at 800°C. (a) As-cast alloy. (b) Extruded alloy (SEM micrographs).
slow oxygen diffusion in g-TiAl and a2-Ti3Al would assist the establishment of a continuous protective alumina-rich film on both phases. As is well known, the high oxygen solubility of a2-Ti3Al [19] or Ti [20,21] alloys and the oxygen dissolution into the metal contribute significantly to the total mass gain in both materials [21–24]. Thus, two different processes give the total oxidation rate: oxide growth and oxygen dissolution into the metal. Over a particular temperature range, one process leads to significantly higher mass gain than does the other. The former process controls the oxidation in pure oxygen at temperatures below 650 and 700°C in pure Ti [20] and Ti–14Al– 21Nb (at.%) alloy [22], respectively, while inward oxygen transport controls oxidation for higher temperatures. In the case of the Ti3Al based alloy, the control of the oxide growth increases up to 800°C when oxidized in air due to the reduction of oxygen entrance by formation of a TiN layer [22]. Although the extruded material presented coarser a2-Ti3Al precipitates, the lower oxygen diffusion would permit the aluminarich layer formed in g-TiAl to undergrow the initially rutile layer developed in a2 regions [14]. This mechanism could be assisted by aluminium diffusion from the adjacent g-TiAl and local Al-enrichment produced by initial rutile formation on a2-Ti3Al lamellae [14].
Fig. 14. Etched cross-sections of (a) as-cast and (b) extruded alloy oxidized for 160 h at 800°C (SEM micrographs).
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Fig. 15. Etched cross-sections of (a) as-cast and (b) extruded alloy oxidized for 160 h at 900°C showing spalling ot the oxide scale (SEM micrographs). Table 1 Composition (in at.%) of the nitride layer, analyzed by EDX, formed on as-cast and extruded alloys after 160 h oxidation at 900°C Alloy
Ti
Al
Mo
N
As-cast Extruded
52.9 48.6
24.1 28.6
0.1 0.2
22.9 22.6
At 800 and 900°C, the thickening of the intermediate alumina layer was not accompanied by a decrease of the oxidation rate as usually observed when alumina controls the oxidation process. The growth of the inner rutile-rich layer and the formation of a new outer rutile layer during further oxidation suggest that the aluminarich layer does not act as an effective barrier reducing the effective surface for inward oxygen or outward titanium diffusion. In addition, the thickening of the intermediate alumina layer would imply outward aluminium diffusion through the inner rutile-rich layer. Etched samples revealed the formation of a continuous nitride layer and an Al-rich a2-free region beneath the nitride layer. These nitrides consisted of TiN and Ti2AlN during the early stages of oxidation, but changed to a Ti2AlN rich layer with increasing exposure time and/or temperature (see Table 1). The estab-
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lishment of a continuous nitride film has been also reported after longer times in Nb-alloyed alloys [6,25,26]. In these Nb-containing alloys, the long time stability of the nitride subsurface layer together with the low diffusivity of cations and anions are considered to be responsible for the excellent resistance of Nb-containing TiAl based alloys [6,26]. The poor spalling resistance of the scale at 900°C was attributed to differences in the thermal expansion coefficients between the oxide scale and the nitride layer. It can be concluded that the presence of the continuous nitride layer at the scale/alloy interface seems to control the oxidation process. The low oxidation rate of titanium nitrides at temperatures below 900°C is well known [27]. Oxidation of this layer leads to a preferential rutile-rich layer. This view is supported by the observed slow growth of the inner rutile-rich layer, and the fact that the fast growth of the rutile-alumina mixed layer reported in binary TiAl [28] or ternary TiAl–Cr [14] did not occur. Thickening of the nitride layer with increasing oxidation time suggests inward nitrogen transport through the oxide scale from the atmosphere to the scale/alloy interface [6,29]. This indicates that the oxide scale does not represent a barrier for inward nitrogen entrance. According to Nickel et al. [6], the high permeability of the alumina-rich barrier is related to rutile-rich inclusions present in this layer. The formation of an Al-rich region, beneath the nitride layer, suggests that dissolution of a2-Ti3Al and Mo-rich phase close to the scale contributed to nitride growth as a supplementary titanium source. Although the presence of this Al-rich zone beneath the scale has not been reported for Mo-containing TiAl based alloys, it has been found in Nb-containing TiAl alloys oxidized in air [6,25,26]. Like the Nb-enrichment observed in the subsurface layer of Nballoyed TiAl alloys, also an enrichment of molybdenum (up to 2 at.%) was detected in this region. The low oxidation rate and the absence of a microstructure influence on the oxidation behaviour of low molybdenum alloyed TiAl alloys indicate the effect of small amounts of this element on the oxidation behaviour of TiAl alloys. This effect will be analyzed considering the applicability of several theories proposed in the literature. It has been reported that additions of certain alloying elements can favour the formation of a protective alumina-rich scale by increasing the Al/Ti activity ratio [30]. This was not the case in the present study, because a thick rutile-rich layer developed at the inner part of the scale rather than an alumina layer. It has been suggested that doping of the outer rutile layer or inner mixed alumina-rutile layer with cations of higher valency than titanium produces a decrease of oxygen vacancy and/or titanium intersti-
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tial concentrations. Thereby, the oxygen diffusion in the outer rutile layer would be reduced, decreasing the rutile growth rate [3,8]. However, the lack of molybdenum in the outer oxide scale proves that this is not the real effect of molybdenum. It has also been proposed that the beneficial effect of niobium additions is related to the direct decrease of anion and cation transport in the inner mixed alumina/rutile layer [31]. In the authors’ opinion, the niobium doping of rutile in the inner part of the scale does not mean a reduction of the oxidation rate because fast oxygen inward transport might occur along the alumina-rutile boundaries. Shida and Anada [32] considered that molybdenum reduces the oxygen solubility in the alloy, causing a transition from internal to external oxidation. These authors associated the beneficial effect of molybdenum to the formation of compact healing alumina layers (an outer layer below the outward rutile scale and inner layer at the scale alloy subsurface). In addition, they simultaneously consider that incorporation of molybdenum into the rutile lattice resulted in a reduction of the defect concentrations in TiO2. The subsequent slow growth of rutile should assist alumina formation. Thus, the establishment of a protective thin alumina barrier at the scale/alloy interface would be favoured because there is more time for the growth or coalescence of alumina particles. Furthermore, these authors reported the formation of a Ti3Al layer beneath the scale at 1000°C, which was not the case in the present study. Rather, an Al-rich subsurface beneath the scale was found in the two Mo-containing alloys. It could be expected that the increased aluminium activity in this sub-surface zone should be followed by the development of an alumina scale. However, in the two alloys studied, the formation of this alumina layer was not found. This suggests that the molybdenum effect, at least in air, was not related to the formation of an inner protective alumina layer. It has been proposed that events occurring at the scale/alloy interface determine the oxidation behaviour of g-TiAl base alloys in pure oxygen and air. In an early stage, a Ti – Al – O single-phase is developed at the Al-depleted zone, which has been designated in the literature as Z-phase [15,28,33], X-phase [29,33], NCP phase [34], Ti2Al [35] or Ti10Al6O2 and Ti10Al6O [36]. The presence of a continuous Ti–Al– O compound in contact with the oxide scale guarantees a slow oxidation rate because alumina is preferentially formed. For longer times or higher temperatures, a considerable increase of the oxidation rate occurs when the single-phase transforms into a two-phase zone, and a2-Ti3Al is lying at the bottom of the scale [37]. The latter results in the formation of a non-protective alumina-rutile mixed
scale, in which alumina formed at Ti–Al–O compound zones and rutile on a2-Ti3Al regions. The high Ti activity, high oxygen solubility and rapid oxygen diffusion should account for the preferential oxidation of a2-Ti3Al to rutile [15,28,38]. However, in the present work, the formation of an Al-rich sub-surface rather than an Al-depleted zone was observed. This indicates that this mechanism is not operating in the two studied Mo-containing alloys. Other authors consider that oxidation is controlled by the formation of a long-term stable nitride layer at the scale alloy interface [6,28,39]. This layer acts as an effective barrier for anion and/or cation transport. This continuous nitride layer has not been observed on binary TiAl based alloys. In this case, the formation of nitrides at a2-Ti3Al regions hampered the formation of an alumina layer at the scale/alloy interface during the initial stages of the oxidation. Nitride formation seems to occur in a similar way for both Mo- or Nb-containing TiAl alloys. However, the development of a continuous nitride layer, whose time for its formation decreases as increases the temperature, indicates that molybdenum affects the behaviour of the g-phase. Molybdenum can decrease the nitrogen solubility and/or the nitrogen diffusivity in the g-phase in such a way that nitride could form when a critical PN2/PO2 ratio is reached at the nitride/alloy interface. Oxygen depletion at the metal–scale interface, due to either oxidation of the alloy or previously formed nitrides, can probably result in an increase of the PN2/PO2 ratio at this region. The latter leads to the formation of a continuous nitride layer. After a continuous nitride layer is established, two processes can take place at the oxide–nitride interface; the oxidation of nitrides and inward diffusion of nitrogen, through the nitride layer, to form new nitride at the scale–alloy interface. The nitride layer would further oxidize to form a rutile-rich layer. Alumina at the inner part of the scale resulted probably from the oxidation of aluminium in Ti2AlN. Scale detachment in both types of microstructures was observed at the highest temperature of 900°C. Spalling commenced after 1 and 10 h at 900°C for the fully lamellar material and duplex material, respectively. According to Becker et al. [40], spalling takes place when a uniform scale reached a critical value (around 10 mm). An identical spalling behaviour has been reported by Pe´rez et al. for a fully-lamellar Ti– 46Al–1Cr–0.2Si alloy, but not for the same alloy with a duplex microstructure [14]. The better adherence of the thicker scale, around 100 mm, of the alloy with duplex microstructure was associated with the formation of a rough scale, with thin and thick regions, more resistant to spalling. This scale should permit an easier
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stress release into the substrate through softer g singlephase regions. In Mo-containing alloys, however, SEM observations prove that scale detachment took place at the oxide–nitride interface (see Fig. 15). This indicates that spalling was provoked by thermal expansion mismatch between the oxide scale and the nitride layer. As the nitride layer is a hard phase, stresses cannot be released through this phase into the alloy, resulting in a similar bad spalling resistance.
5. Conclusions From the present study, it can be concluded that the oxidation behaviour of Ti – 46.8Al – 1Mo – 0.2Si alloy in air does not depend on the type of microstructure, i.e. lamellar or duplex. The low oxidation rate can be attributed to a combined effect of a compact intermediate alumina-rich layer that decreases inward oxygen diffusion and favours the formation of a continuous layer of titanium nitride at the scale – alloy interface. The molybdenum effect could be related with the decrease of nitrogen solubility and/or nitrogen diffusivity in the g-phase. The thickening of the nitride layer with progressing oxidation arises from inward nitrogen transport from the atmosphere to the scale – alloy interface. The formation of the nitride layer avoids preferential oxidation through a2-Ti3Al phase, thus preventing the development of a non-protective inner alumina-rutile mixed layer.
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[10] A. Gil, H. Hoven, E. Wallura, W.J. Quadakkers, Corros. Sci. 34 (1993) 615. [11] M.P. Brady, J.L. Smialek, E.D. Verink Jr, D.T. Hoelzer, R. Stone, Mat. Manuf. Processes 11 (1996) 635. [12] V.A.C. Haanappel, R. Hofman, J.D. Sunderko¨tter, W. Glatz, H. Clemens, M.F. Stroosnijder, Oxid. Met. 48 (1997) 263. [13] V.A.C. Haanappel, W. Glatz, H. Clemens, M.F. Stroosnijder, Mat. High Temp. 14 (1997) 19. [14] P. Pe´rez, J.A. Jime´nez, G. Frommeyer, P. Adeva, Oxid Met. 53 (2000) 99. [15] V. Shemet, H. Hoven, W.J. Quadakkers, Intermetallics 5 (1997) 311. [16] H. Clemens, A. Lorich, N. Eberhardt, W. Glatz, W. Knabl, H. Kestler, Z. Metallkd. 90 (1999) 569. [17] S.L. Semiatin, N. Frey, S.M. El-Soudani, J.D. Bryant, Metall. Trans. A 23A (1994) 1719. [18] C. Koeppe, A. Bartels, J. Seeger, H. Mecking, Metall. Trans. A 24A (1993) 1795. [19] A. Menand, A Huguet, A. Nerac-Partaix, Acta Mater. 44 (1996) 4729. [20] J. Stringer, Acta Metall. 8 (1960) 758. [21] P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, London, 1988, p. 292. [22] T.A. Wallace, R.K. Clark, K.E. Wiedemann, Oxid. Met. 42 (1994) 451. [23] R.U. Vaidya, Y.S. Park, J. Zhe, G.T. Gray III, D.P. Butt, Oxid. Met. 50 (1998) 215. [24] J. Unnam, R.N. Shenoy, R.K. Clark, Oxid. Met. 26 (1986) 231. [25] I.E. Locci, M.P. Brady, R.A. MacKay, J.W. Smith, Scripta Mater. 37 (1997) 761. [26] U. Figge, A. Elschner, N. Zheng, H. Schuster, W.J. Quadakkers, Fresenius J. Anal. Chem. 346 (1993) 75. [27] J. Desmaison, P. Lefort, M. Billy, Oxid. Met. 13 (1979) 203. [28] N. Zheng, W.J. Quadakkers, A. Gil, H. Nickel, Oxid. Met. 44 (1995) 477. [29] F. Dettenwanger, E. Schumann, M. Ru¨hle, J. Rakowski, G.H. Meier, Oxid. Met. 50 (1998) 269. [30] S. Becker, A. Rahmel, M. Schorr, M. Schu¨tze, Oxid. Met. 38 (1992) 425. [31] M.F. Stroosnijder, N. Zheng, W.J. Quadakkers, R. Hofman, A. Gil, F. Lanza, Oxid. Met. 46 (1996) 19. [32] Y. Shida, H. Anada, Oxid. Met. 45 (1996) 197. [33] M. Groß, V. Kolarik, A. Rahmel, Oxid. Met. 48 (1997) 171. [34] C. Lang, M. Schu¨tze, Mater. Corros. 48 (1997) 13. [35] W.E. Dowling Jr, W.T. Donlon, Scrip. Metall. Mater. 27 (1992) 1663. [36] R.W. Beye, R. Gronsky, Acta Metall. Mater. 42 (1994) 1373. [37] W.J. Quadakkers, N. Zheng, A. Gil, E. Wallura, H. Hoven, Mater. Sci. Forum 251-254 (1997) 187. [38] N. Zheng, W. Fischer, H. Gru¨bmeier, V. Shemet, W.J. Quadakkers, Scrip. Metall. Mater. 33 (1995) 47. [39] W.J. Quadakkers, P. Schaaf, N. Zheng, A. Gil, E. Wallura, Mater. Corros. 48 (1997) 28. [40] S. Becker, M. Schu¨tze, A. Rahmel, Oxid. Met. 39 (1993) 93.
Oxidation behaviour of a Ti–46Al–1Mo–0.2Si alloy: the effect of Mo addition and alloy microstructure P. Pe´rez a,b,*, J.A. Jime´nez b, G. Frommeyer c, P. Adeva b b
a Institute for Health and Consumer Protection. JRC of the European Commision, I-21020 (Va), Ispra, Italy Centro Nacional de In6estigaciones Metalu´rgicas, CENIM, (CSIC). A6da. Gregorio del Amo 8, 28040, Madrid, Spain c Max-Planck-Institut fu¨r Eisenforschung GmbH, Max-Planck-Strabe 1, D-40237 Du¨sseldorf, Germany
Received 14 April 1999; received in revised form 21 January 2000
Abstract The influence of molybdenum addition and alloy microstructure on the oxidation behaviour of a Ti – 46.8Al – 1Mo –0.2Si alloy was studied in air at temperatures ranging from 600 to 900°C. The alloy produced by arc melting exhibited a structure of coarse lamellar grains in as-cast condition that transformed to a duplex microstructure after hot extrusion at 1300°C. Oxidation rate and scale spallation resistance were not affected by the type of microstructure. The effect of molybdenum addition was related to the formation and thickening of a protective continuous nitride layer. The low mass gain measured can be associated with the low oxidation rate of the nitride layer. Finally, the poor spalling resistance of the oxide scale at 900°C was attributed to the mismatch in thermal expansion coefficients of the oxide scale and the nitride layer. © 2000 Elsevier Science S.A. All rights reserved. Keywords: g-TiAl based intermetallic alloy; High temperature oxidation; Microstructure; Alloying
1. Introduction The research of structural materials for high-temperature applications has led to the consideration of gTiAl base alloys for aviation gas turbines or automotive engine components, since this intermetallic compound combines low density with high specific strength and stiffness up to 800°C. However, low ductility at room temperature, poor fracture toughness and insufficient oxidation resistance at temperatures above 700°C have limited the use of g-TiAl based alloys. Attempts for improving these properties have been made through additions of specific alloying elements and microstructural modifications. Several investigations have shown that V, Mn, and Cr generally increase the ductility [1], and elements like Mo, Nb, Hf, Zr, Cl or W improve the oxidation behaviour [2 – 8]. However, each of these alloying elements improves a specific property but can be detrimental to the other properties. Therefore, it is difficult to balance good mechanical properties and oxidation resistance only by alloying elements. * Corresponding author.
Changes in the microstructural morphology have been also reported to affect strongly the mechanical properties of g-based alloys [1]. For two-phase g-TiAl alloys, different microstructures can be generated depending on processing route and/or subsequent thermal treatments, i.e. fully lamellar, nearly lamellar, duplex or near g [1]. Generally, the duplex structure is desirable for ductility, but poor for toughness and creep resistance. The opposite is true for the fully lamellar structure. At this point, an adequate ternary element addition and control of alloy microstructure could be appropriated to improve both, mechanical properties and oxidation resistance of these alloys. However, the effect of type of microstructure on the oxidation behaviour has not been thoroughly investigated, as indicated by the controversy existing in the literature about this point [9–13]. Gil et al. [9,10] reported for a two-phase g/a2TiAl alloy oxidized at 900°C in air, that the oxidation rate and type of the formed scale depend on the type of microstructure and distribution of both phases. Cast material with a fine lamellar microstructure formed a thin protective scale instead of the thicker rutile-based scale established on heat treated material, whose mi-
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crostructure consisted of large grains of g-TiAl and coarse a2-Ti3Al. Recently, Pe´rez et al. [14] found important differences in the oxidation behaviour of fullylamellar and duplex Ti – 46Al – 1Cr – 0.2Si alloy. At temperatures ranging from 700 to 900°C, the duplex microstructure presented a higher oxidation rate than the fully-lamellar. On the other hand, Shemet et al. [15], for TiAl oxidized in argon/oxygen at 900°C and Haanappel et al. [12,13] for Ti – 48Al – 2Cr oxidized in air at 700 and 800°C, reported that oxidation rate was unaffected by the alloy microstructure. The aim of this study is to determine the influence of the type of microstructure, i.e. fully lamellar in as-cast condition and duplex after extrusion, on the oxidation behaviour of TiAl based alloys containing small additions of Mo. This work is complementary to a previous one on the influence of the microstructure on the oxidation behaviour of alloys containing small amounts of Cr (1 at.%), which is detrimental for the oxidation resistance [14]. In that case, it was found a considerable effect of the alloy microstructure on the oxidation behaviour. In the present work, the oxidation behaviour of fully-lamellar and duplex Ti – 46.8Al–1Mo– 0.2Si alloy was studied. Oxidation kinetics have been evaluated in air at temperatures ranging from 600 to 900°C.
2. Experimental A Mo-alloyed two-phase g/a2-TiAl alloy with composition (in at.%) Ti – 46.8Al – 1Mo – 0.19Si was produced in 50 kg ingots by arc skull melting. As-cast ingots exhibited a coarse-grained microstructure. This microstructure was refined by hot extrusion of billets 75 mm in diameter and 120 mm in length in the a + g field at 1300°C. The reduction in area of the extruded rods was about 7:1. Microstructure present in the as-cast material and after extrusion was characterized using X-ray diffraction, optical and scanning electron microscopy (SEM). The phases present and the volume fraction of each phase were determined by X-ray diffraction. The diffraction patterns showed the presence of reflections of both phases g-TiAl and a2-Ti3Al. The volume fraction of each phase was calculated from the integrated intensities of (201), (022), (220), and (203) peaks of the a2-Ti3Al, and (002), (200), (202), and (220) peaks of g-TiAl. A decrease in the amount of the a2-Ti3Al from 20 vol.% in the as-cast material to 10 vol.% after hot extrusion was found. This value agrees with that predicted by the equilibrium diagram, about 10 vol.% [16]. Oxidation tests were carried out on rectangular coupons of 10× 10×1 mm3 cut from the billets by electrospark erosion. The surfaces were abraded on successively finer silicon carbide papers, then mechani-
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cally polished with 1 mm diamond paste and finally ultrasonically cleaned with acetone. Oxidation kinetics in the temperature range from 600 to 800°C were determined by discontinuous thermogravimetry. The actual test temperature was verified by a thermocouple placed close to the sample. After exposure, samples were taken out of the furnace and left to cool down, then weighed using a balance with a resolution of 910 mg. Characterization of oxidation products was performed on samples isothermally oxidized for short and intermediate exposure times, as well as on samples used for the oxidation kinetics determination. Cross-sections of samples were prepared by conventional metallographic technique. To prevent scale loss during metallographic preparation sample surfaces were successively coated with a thin gold layer (by sputtering) and a thicker layer of copper (electrolytically deposited). Cross-sections were prepared by conventional metallographical techniques. Surfaces and cross-sections of the oxidized specimens were studied by SEM. Phase identification of the oxide scale was performed by X-ray diffraction (XRD), energy dispersive X-ray microanalysis (EDX) and wavelength dispersion X-ray microanalysis (WDX).
3. Results
3.1. Microstructure As-cast alloy presented a microstructure consisting of very coarse lamellar grains with a size of several hundred micrometers, as shown in Fig. 1(a). This lamellar structure was partially transformed by extrusion at 1300°C into a duplex microstructure consisting of single-phase g grains, ranging from 5 to 20 mm in size, and finer grains, as shown in Fig. 1(b). Observations in the extrusion direction showed that both single-phase g and two-phase g/a2 grains were grouped in alternating bands ranging from 60 to 150 mm in width (Fig. 2). The origin of such microstructure has been attributed to the segregation of aluminium during the peritectic solidification [17,18]. However, in the present work the development of this stripped microstructure must be related to the precipitation of small amounts of a Mo-rich phase, as confirmed by EDX microanalysis. These Morich particles, probably Ti2AlMo as found by XRD (see Fig. 3), are heterogeneously distributed on the microstructure, as observed in Fig. 4(a and b) corresponding to bands of coarse and fine grain, respectively. Mo-rich precipitates were only present in fine grain regions (see Fig. 4(b)). This non-homogeneous distribution results from segregation of molybdenum during the casting process.
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As Mo-rich particles were not found in the as-cast material, it was concluded that they precipitated before the extrusion, during heating. In the extrusion process, these particles pinned the grain boundaries of the recrystallized grains, preventing the grain growth during and after extrusion. However, as distribution of these
Fig. 3. XRD spectra of non-oxidized extruded alloy.
Fig. 1. Optical micrographs showing the microstructure of (a) as-cast material and (b) extruded material.
Fig. 4. Micrographs of the microstructure in coarse (a) and fine grain regions (b) of the extruded alloy. It is noticed the presence of Mo-rich precipitates exclusively at grain boundaries of fine grain regions.
Fig. 2. SEM micrograph of the extruded alloy showing the microstructure in the extrusion direction.
particles is heterogeneous, in those regions with a low volume fraction of particles the grain growth will not be inhibited, resulting in the stripped microstructure.
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3.2. Oxidation kinetics Oxidation kinetics of the Ti – 46.8Al – 1Mo –0.2Si alloy have been evaluated in air at temperatures ranging from 600 to 900°C by discontinuous thermogravimetry. The mass gain– time curves for both extruded and as-cast material are represented in Fig. 5. The general oxidation behaviour of the two microstructures, i.e. duplex and lamellar, is quite similar in the whole temperature range tested. At 600°C, the mass gains were negligible and the curves for both materials can be superimposed. At 700 and 800°C, the oxidation rate of both as-cast and extruded alloy tended to decrease with progressing oxidation until a steady state was reached. At 900°C, the kinetics were affected by an extensive scale spallation already after short exposure times, 1 and 10 h for as-cast and extruded material, respectively.
Fig. 6. Surface morphology of the scale formed after 160 h of oxidation at 600°C, showing microstructure of (a) as-cast and (b) extruded alloy (SEM micrographs).
3.3. Surface morphology
Fig. 5. Mass gain-time curves for both as-cast and extruded alloy oxidized in air at: (a) 600 and 700°C. (b) 800 and 900°C.
The thin scale formed after testing at 600°C revealed the microstructure of the alloy because dark and bright zones were formed above the g-TiAl and a2-Ti3Al phases, respectively (see Fig. 6). EDX analysis evidenced that the dark zones consisted of an Al-rich oxide, likely alumina, while the bright zones consisted of an Al-rich oxide covered by a very porous and discontinuous Ti-rich oxide, probably rutile. At 700°C, an alumina-rich layer formed at the surface from the early stages of oxidation. For longer exposure times, the rutile content of the scale increases and small well-developed rutile crystals commenced to appear progressively on the surface, as shown in Fig. 7. No differences between as-cast and extruded material were observed. At 800 and 900°C, the morphology of the scale was completely different from that described before. The outer layer consisted of the rutile phase, even in the earlier stages of oxidation. The size of the rutile crystals increases with increasing temperature and exposure time, as shown in Fig. 8.
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3.4. Cross-sectional obser6ations
Fig. 7. SEM micrographs showing surface scale morphology after 150 h at 700°C; (a) as-cast alloy, (b) extruded alloy.
Cross sectional examination revealed only small differences between the scale formed on the as-cast and extruded material. A thin uniform scale developed on both materials in the whole temperature range studied, as shown in Fig. 9. At 600°C, this scale was too thin to be analyzed. At 700°C, the scale formed after 50 h of exposure was very thin (B 1 mm). It consisted of two layers, an outer alumina-rich layer and an inner rutilerich layer. Low scale growth rate was observed during further oxidation. Thus, after 150 h of exposure an oxide scale 3 mm thick was observed, in which the inner part has a thickness of 2 mm and the outer part a thickness of 1 mm (see Fig. 10). A nearly continuous white Ti-rich phase appeared at the scale/alloy interface. At 800 and 900°C, the scale pattern for short oxidation times was practically identical to that observed at 700°C for longer times. The major difference concerned the development of a thin outer rutile layer (Fig. 11). The Ti-rich phase at the scale/alloy interface was identified by XRD as titanium nitrides of the type Ti2AlN (Fig. 12). In addition, also very small peaks of TiN (not observed in Fig. 12) were detected for both microstructures. After 160 h at 800°C, the scale was 10 mm thick and all the layers composing the scale thickened (Fig. 13). The intermediate alumina-rich layer was 2 mm thick and presented an apparently dense microstructure. Nevertheless, this alumina layer was not sufficiently protective because it prevented neither the
Fig. 8. SEM micrographs of the oxide scale morphology developed on as-cast (a, b) and extruded alloy (c, d) after 160 h of exposure at 800°C (a, c) and 900°C (b, d).
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exposure times at 900°C (Fig. 15). EDX analysis suggested that nitride layer was composed mainly of Ti2AlN (see Table 1). In addition, it was checked that sub-scale region contained a small amount of oxygen, about 3%.
4. Discussion The results obtained in the present work show that the mechanism controlling the oxidation of the Ti– 46Al–1Mo–0.2 Si alloy was very similar for both microstructures, i.e. fully-lamellar and duplex. The oxidation rate at 600, 700, 800 and 900°C was practically identical for both microstructures. At 600°C, the low oxidation rate suggests the formation of a continuous healing alumina-rich film over the entire surface. Above the alumina-rich layer situated at a2-Ti3Al lamellae appeared a porous rutile layer, probably formed during the initial stages of oxidation. The
Fig. 9. Cross-sections of the scales formed on (a, c, e) as-cast and (b, d, f) extruded alloy after 150 h 700°C (a, b), 160 h 800°C (c, d) and 1 h 900 (e, f) (SEM micrographs).
Fig. 11. Cross-section of the scale formed after 10 h at 800°C on the extruded alloy (SEM micrograph).
Fig. 10. Cross-sections of the scale formed after 150 h at 700°C on the extruded alloy (SEM micrograph).
outward diffusion of titanium nor the inward diffusion of oxygen. This resulted in a slow but continuous growth of the outermost and inner rutile-rich layers. The coalescence of the innermost titanium nitrides led to the formation of a continuous layer, 1 mm in thickness, at the scale/alloy interface. A SEM micrograph of the sample presented in Fig. 13 after etching is shown in Fig. 14. This figure revealed the formation of an Al-rich a2-Ti3Al free zone beneath the nitride layer. Both, the Al-rich a2-Ti3Al free zone and the nitride layer significantly thickened after longer
Fig. 12. XRD pattern of oxide scale formed on the extruded alloy after 24 h of oxidation at 800°C.
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At 700°C, the low oxidation rate measured in the early stages of oxidation for both alloys suggests that an alumina-rich film formed in this period. As oxidation proceeds, surface morphology observations showed the presence of rutile crystals, indicative of outward titanium transport through the scale. The appearance of these rutile crystals seems to account for the increase of the oxidation rate measured in the mass gain curves. At this stage, the alumina-rich scale left to be protective as proves the development of an inner thicker rutilerich layer. The formation of this layer implies inward oxygen through the outer alumina-rich layer. The decrease of the oxidation rate for longer times suggests that a more protective layer started to control the oxidation. It could be possible that thickening of the intermediate alumina-rich layer with time leads to a more compact and protective scale. Nevertheless, one might assume also that a new more protective layer starts to control the oxidation process. This new layer could correspond with the nearly continuous titanium nitride layer detected at the scale/alloy interface after longer times. The formation of a continuous nitride layer and further thickening should occur for long term exposures as it has been described for the oxidation of Ti–48Al–2Cr–2Nb alloy in air at 700°C [25]. Fig. 13. Cross-sections of the scales formed after 160 h of exposure at 800°C. (a) As-cast alloy. (b) Extruded alloy (SEM micrographs).
slow oxygen diffusion in g-TiAl and a2-Ti3Al would assist the establishment of a continuous protective alumina-rich film on both phases. As is well known, the high oxygen solubility of a2-Ti3Al [19] or Ti [20,21] alloys and the oxygen dissolution into the metal contribute significantly to the total mass gain in both materials [21–24]. Thus, two different processes give the total oxidation rate: oxide growth and oxygen dissolution into the metal. Over a particular temperature range, one process leads to significantly higher mass gain than does the other. The former process controls the oxidation in pure oxygen at temperatures below 650 and 700°C in pure Ti [20] and Ti–14Al– 21Nb (at.%) alloy [22], respectively, while inward oxygen transport controls oxidation for higher temperatures. In the case of the Ti3Al based alloy, the control of the oxide growth increases up to 800°C when oxidized in air due to the reduction of oxygen entrance by formation of a TiN layer [22]. Although the extruded material presented coarser a2-Ti3Al precipitates, the lower oxygen diffusion would permit the aluminarich layer formed in g-TiAl to undergrow the initially rutile layer developed in a2 regions [14]. This mechanism could be assisted by aluminium diffusion from the adjacent g-TiAl and local Al-enrichment produced by initial rutile formation on a2-Ti3Al lamellae [14].
Fig. 14. Etched cross-sections of (a) as-cast and (b) extruded alloy oxidized for 160 h at 800°C (SEM micrographs).
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Fig. 15. Etched cross-sections of (a) as-cast and (b) extruded alloy oxidized for 160 h at 900°C showing spalling ot the oxide scale (SEM micrographs). Table 1 Composition (in at.%) of the nitride layer, analyzed by EDX, formed on as-cast and extruded alloys after 160 h oxidation at 900°C Alloy
Ti
Al
Mo
N
As-cast Extruded
52.9 48.6
24.1 28.6
0.1 0.2
22.9 22.6
At 800 and 900°C, the thickening of the intermediate alumina layer was not accompanied by a decrease of the oxidation rate as usually observed when alumina controls the oxidation process. The growth of the inner rutile-rich layer and the formation of a new outer rutile layer during further oxidation suggest that the aluminarich layer does not act as an effective barrier reducing the effective surface for inward oxygen or outward titanium diffusion. In addition, the thickening of the intermediate alumina layer would imply outward aluminium diffusion through the inner rutile-rich layer. Etched samples revealed the formation of a continuous nitride layer and an Al-rich a2-free region beneath the nitride layer. These nitrides consisted of TiN and Ti2AlN during the early stages of oxidation, but changed to a Ti2AlN rich layer with increasing exposure time and/or temperature (see Table 1). The estab-
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lishment of a continuous nitride film has been also reported after longer times in Nb-alloyed alloys [6,25,26]. In these Nb-containing alloys, the long time stability of the nitride subsurface layer together with the low diffusivity of cations and anions are considered to be responsible for the excellent resistance of Nb-containing TiAl based alloys [6,26]. The poor spalling resistance of the scale at 900°C was attributed to differences in the thermal expansion coefficients between the oxide scale and the nitride layer. It can be concluded that the presence of the continuous nitride layer at the scale/alloy interface seems to control the oxidation process. The low oxidation rate of titanium nitrides at temperatures below 900°C is well known [27]. Oxidation of this layer leads to a preferential rutile-rich layer. This view is supported by the observed slow growth of the inner rutile-rich layer, and the fact that the fast growth of the rutile-alumina mixed layer reported in binary TiAl [28] or ternary TiAl–Cr [14] did not occur. Thickening of the nitride layer with increasing oxidation time suggests inward nitrogen transport through the oxide scale from the atmosphere to the scale/alloy interface [6,29]. This indicates that the oxide scale does not represent a barrier for inward nitrogen entrance. According to Nickel et al. [6], the high permeability of the alumina-rich barrier is related to rutile-rich inclusions present in this layer. The formation of an Al-rich region, beneath the nitride layer, suggests that dissolution of a2-Ti3Al and Mo-rich phase close to the scale contributed to nitride growth as a supplementary titanium source. Although the presence of this Al-rich zone beneath the scale has not been reported for Mo-containing TiAl based alloys, it has been found in Nb-containing TiAl alloys oxidized in air [6,25,26]. Like the Nb-enrichment observed in the subsurface layer of Nballoyed TiAl alloys, also an enrichment of molybdenum (up to 2 at.%) was detected in this region. The low oxidation rate and the absence of a microstructure influence on the oxidation behaviour of low molybdenum alloyed TiAl alloys indicate the effect of small amounts of this element on the oxidation behaviour of TiAl alloys. This effect will be analyzed considering the applicability of several theories proposed in the literature. It has been reported that additions of certain alloying elements can favour the formation of a protective alumina-rich scale by increasing the Al/Ti activity ratio [30]. This was not the case in the present study, because a thick rutile-rich layer developed at the inner part of the scale rather than an alumina layer. It has been suggested that doping of the outer rutile layer or inner mixed alumina-rutile layer with cations of higher valency than titanium produces a decrease of oxygen vacancy and/or titanium intersti-
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tial concentrations. Thereby, the oxygen diffusion in the outer rutile layer would be reduced, decreasing the rutile growth rate [3,8]. However, the lack of molybdenum in the outer oxide scale proves that this is not the real effect of molybdenum. It has also been proposed that the beneficial effect of niobium additions is related to the direct decrease of anion and cation transport in the inner mixed alumina/rutile layer [31]. In the authors’ opinion, the niobium doping of rutile in the inner part of the scale does not mean a reduction of the oxidation rate because fast oxygen inward transport might occur along the alumina-rutile boundaries. Shida and Anada [32] considered that molybdenum reduces the oxygen solubility in the alloy, causing a transition from internal to external oxidation. These authors associated the beneficial effect of molybdenum to the formation of compact healing alumina layers (an outer layer below the outward rutile scale and inner layer at the scale alloy subsurface). In addition, they simultaneously consider that incorporation of molybdenum into the rutile lattice resulted in a reduction of the defect concentrations in TiO2. The subsequent slow growth of rutile should assist alumina formation. Thus, the establishment of a protective thin alumina barrier at the scale/alloy interface would be favoured because there is more time for the growth or coalescence of alumina particles. Furthermore, these authors reported the formation of a Ti3Al layer beneath the scale at 1000°C, which was not the case in the present study. Rather, an Al-rich subsurface beneath the scale was found in the two Mo-containing alloys. It could be expected that the increased aluminium activity in this sub-surface zone should be followed by the development of an alumina scale. However, in the two alloys studied, the formation of this alumina layer was not found. This suggests that the molybdenum effect, at least in air, was not related to the formation of an inner protective alumina layer. It has been proposed that events occurring at the scale/alloy interface determine the oxidation behaviour of g-TiAl base alloys in pure oxygen and air. In an early stage, a Ti – Al – O single-phase is developed at the Al-depleted zone, which has been designated in the literature as Z-phase [15,28,33], X-phase [29,33], NCP phase [34], Ti2Al [35] or Ti10Al6O2 and Ti10Al6O [36]. The presence of a continuous Ti–Al– O compound in contact with the oxide scale guarantees a slow oxidation rate because alumina is preferentially formed. For longer times or higher temperatures, a considerable increase of the oxidation rate occurs when the single-phase transforms into a two-phase zone, and a2-Ti3Al is lying at the bottom of the scale [37]. The latter results in the formation of a non-protective alumina-rutile mixed
scale, in which alumina formed at Ti–Al–O compound zones and rutile on a2-Ti3Al regions. The high Ti activity, high oxygen solubility and rapid oxygen diffusion should account for the preferential oxidation of a2-Ti3Al to rutile [15,28,38]. However, in the present work, the formation of an Al-rich sub-surface rather than an Al-depleted zone was observed. This indicates that this mechanism is not operating in the two studied Mo-containing alloys. Other authors consider that oxidation is controlled by the formation of a long-term stable nitride layer at the scale alloy interface [6,28,39]. This layer acts as an effective barrier for anion and/or cation transport. This continuous nitride layer has not been observed on binary TiAl based alloys. In this case, the formation of nitrides at a2-Ti3Al regions hampered the formation of an alumina layer at the scale/alloy interface during the initial stages of the oxidation. Nitride formation seems to occur in a similar way for both Mo- or Nb-containing TiAl alloys. However, the development of a continuous nitride layer, whose time for its formation decreases as increases the temperature, indicates that molybdenum affects the behaviour of the g-phase. Molybdenum can decrease the nitrogen solubility and/or the nitrogen diffusivity in the g-phase in such a way that nitride could form when a critical PN2/PO2 ratio is reached at the nitride/alloy interface. Oxygen depletion at the metal–scale interface, due to either oxidation of the alloy or previously formed nitrides, can probably result in an increase of the PN2/PO2 ratio at this region. The latter leads to the formation of a continuous nitride layer. After a continuous nitride layer is established, two processes can take place at the oxide–nitride interface; the oxidation of nitrides and inward diffusion of nitrogen, through the nitride layer, to form new nitride at the scale–alloy interface. The nitride layer would further oxidize to form a rutile-rich layer. Alumina at the inner part of the scale resulted probably from the oxidation of aluminium in Ti2AlN. Scale detachment in both types of microstructures was observed at the highest temperature of 900°C. Spalling commenced after 1 and 10 h at 900°C for the fully lamellar material and duplex material, respectively. According to Becker et al. [40], spalling takes place when a uniform scale reached a critical value (around 10 mm). An identical spalling behaviour has been reported by Pe´rez et al. for a fully-lamellar Ti– 46Al–1Cr–0.2Si alloy, but not for the same alloy with a duplex microstructure [14]. The better adherence of the thicker scale, around 100 mm, of the alloy with duplex microstructure was associated with the formation of a rough scale, with thin and thick regions, more resistant to spalling. This scale should permit an easier
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stress release into the substrate through softer g singlephase regions. In Mo-containing alloys, however, SEM observations prove that scale detachment took place at the oxide–nitride interface (see Fig. 15). This indicates that spalling was provoked by thermal expansion mismatch between the oxide scale and the nitride layer. As the nitride layer is a hard phase, stresses cannot be released through this phase into the alloy, resulting in a similar bad spalling resistance.
5. Conclusions From the present study, it can be concluded that the oxidation behaviour of Ti – 46.8Al – 1Mo – 0.2Si alloy in air does not depend on the type of microstructure, i.e. lamellar or duplex. The low oxidation rate can be attributed to a combined effect of a compact intermediate alumina-rich layer that decreases inward oxygen diffusion and favours the formation of a continuous layer of titanium nitride at the scale – alloy interface. The molybdenum effect could be related with the decrease of nitrogen solubility and/or nitrogen diffusivity in the g-phase. The thickening of the nitride layer with progressing oxidation arises from inward nitrogen transport from the atmosphere to the scale – alloy interface. The formation of the nitride layer avoids preferential oxidation through a2-Ti3Al phase, thus preventing the development of a non-protective inner alumina-rutile mixed layer.
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