- Email: [email protected]
Influence of water vapour on the oxidation behaviour of titanium aluminides
Intermetallics 10 (2002) 59–72 www.elsevier.com/locate/intermet
Influence of water vapour on the oxidation behaviour of titanium aluminides A. Zeller, F. Dettenwanger, M. Schu¨tze* Karl-Winnacker-Institut der DECHEMA e.V., Theodor-Heuss-Allee 25, D-60486 Frankfurt am Main, Germany Received 29 May 2001; accepted 15 August 2001
Abstract The influence of water vapour on the oxidation of a technical g-TiAl based alloy at temperatures of 700 and 750 C is discussed. The present paper focuses on the oxidation behaviour in dry air and air containing 10 vol.% H2O with respect to oxidation kinetics and microstructural investigations of the oxide scales and the subsurface zones. The presence of water vapour led to a significant acceleration in the oxidation kinetics as well as to a change in the crystal morphology and microstructure of the oxide scales formed. The results indicate that the kinetics of the reaction of water molecules with the rutile phase and a change in its defect structure are important parameters in the mechanism. The subsurface zone formed in water vapour-containing atmospheres consisted (in addition to the cubic Al-depleted layer) of a lamellar layer of Al2O3 and Ti due to either internal oxidation or decomposition of the cubic phase. Both, the oxide scale and the depletion layer are important regarding the mechanical behaviour of a TiAl component at high temperatures which was also investigated. # 2002 Elsevier Science Ltd. All rights reserved. Keywords: A. Titanium aluminides based on TiAl; B. Oxidation; B. Phase identification; B. Environmental embrittlement
1. Introduction Titanium aluminides based on the intermetallic phase g-TiAl are attractive as a new group of high temperature alloys due to their specific high temperature strength. Research undertaken in the past 10 years was mainly concentrated on the understanding and improvement of the mechanical properties and oxidation behaviour of these alloys. However, the influence of water vapour on the oxidation behaviour and the mechanical properties of titanium aluminides has not been studied in the literature to a great extent yet although in the envisaged applications as gas turbine blades or car engine valves the atmosphere will contain a significant amount of H2O (up to 20 vol.%). In a more general sense water vapour can affect several properties of a high temperature component through the influence on the formation, microstructure, mechanical properties, adhesion, and transport properties of oxide scales [1–4] as well as the mechanical properties of the base metal directly. However, it has to be kept in mind that the influence of water vapour on the high * Corresponding author. Tel.: +49-69-7564-361; fax: +49-69-7564368.
temperature properties of alloys strongly depends on the specific system being considered, which makes a general modelling difficult if not impossible. A systematic study concerning the influence of water vapour on the oxidation of g-TiAl based alloys was conducted by Kremer et al. at 900 C in pure oxygen and oxygen with various water vapour contents [5,6]. The presence of water vapour led to a significant increase in the oxidation rate compared to the oxidation in dry oxygen. Interestingly, the accelerated kinetics was connected with an incubation time and the mass increase during the early stage was less in water vapour-containing atmosphere compared to dry conditions. After longer times ( 50 h) a linear oxidation rate was observed which was slowed down by switching the atmosphere to dry conditions. Changing back to the water vapourcontaining atmosphere accelerated the oxidation rate again after an incubation period. Variation of the water vapour and oxygen partial pressures revealed a variation of the linear rate constant with pH2O1/2 and pO21/4 which was interpreted with respect to different dissociation kinetics of the water molecule compared to the oxygen molecule. The different kinetics were reflected in the microstructures of the oxide scales formed. In dry oxygen
0966-9795/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(01)00104-2
60
A. Zeller et al. / Intermetallics 10 (2002) 59–72
a dense protective alumina barrier was formed whereas in humid oxygen only islands of alumina were present in the outer TiO2 layer. A possible mechanism of doping the oxide by hydrogen produced due to the dissociation of H2O was proposed [6]. Brady et al. [7,8] compared the oxidation behaviour of several TiAl,Cr-alloys with focus on the Cr-content at 1000 C in oxygen, and dry and wet air. In oxygen a dense alumina scale was observed on all alloys. In air the alloys with high Cr contents (15 at.%) consisting of g-TiAl and the Laves phase Ti(Cr,Al)2 formed a dense alumina scale independent of either dry or humidified air. This formation was promoted by the presence of the Laves phase. Therefore, an influence of water vapour was not observed for these alloys. In contrast, g-TiAl alloys with low Cr contents (< 5 at.%) showed a strong acceleration of the oxidation kinetics for the moist air. In addition, the structure of the subsurface zone could be correlated with the exposure conditions. The formation of a b-Ti(Al,Cr)-phase at the metal/oxide phase boundary hindering the formation of a protective alumina layer was observed for the moist atmosphere. It was concluded that the composition of the phase(s) present at the metal/scale interface is directly related to the influence of the water vapour in the atmosphere on the oxidation kinetics, i.e., on the formation of a protective alumina scale. Hald [9] observed an acceleration of the oxidation kinetics of TiAl at 900 C by switching from dry to moist air for the case of a non protective oxide scale on TiAl. For a previously formed alumina scale the switch to moist air did not lead to accelerated oxidation kinetics. Takasaki et al. [10] studied the influence of hydrogen dissolution in two phase titanium aluminides on the oxidation behaviour at 897 and 1077 C. The time of hydrogen loading influenced the subsequent oxidation behaviour significantly. Short loading times and therefore small amounts of dissolved hydrogen led to smaller weight gains compared to unloaded TiAl. Longer loading times led to higher weight gains which surprisingly increased with increasing Al content (Ti42Al, Ti45Al, Ti50Al). The beneficial effect of small hydrogen contents was attributed to a reduction of the Ti diffusion in the alloy, an explanation for the dependence of the Al content after long hydrogen loading times was not given. Investigations concerning the influence of water vapour on the mechanical and oxidation properties of gTiAl based alloys at the potential application temperatures of 700–750 C are to the authors’ knowledge not present in the literature. We investigated these points for a technical g-TiAl based alloy at temperatures of 700 and 750 C and the results will be presented in two papers. The present paper concentrates on the influence of water vapour on the oxidation behaviour and oxide scale microstructure for scales developed on the alloy with and without mechanical loading.
A second paper concentrates on the influence of water vapour on the creep and fatigue properties of the alloy Ti 47Al 1Cr Si and has been submitted to this journal.
2. Experimental The alloy Ti 47Al 1Cr Si investigated was provided by Tital, Bestwig, Germany, in the form of near specimen shaped slugs for the mechanical testing specimens in the as cast condition. The chemical composition is given in Table 1 and is close to the composition of a technical gbased TiAl alloy developed for application in automotive engines. The microstructure of the alloy is described in [11]. The material was cast in the near net shape of the cylindrical creep specimens. Fig. 1 illustrates geometry and locations of the oxidation and creep specimens cut from the as-received ingot. The oxidation specimens were cut from the lower part of the ingot with the dimensions 20102 mm3. The final grinding step for the coupons was performed with a 500 grit SiC paper and after grinding the coupons were ultrasonically cleaned in ethanol and acetone prior to the oxidation tests. The oxidation coupons were used for thermogravimetric measurements carried out in synthetic air and synthetic air containing 10 vol.% H2O at 700 and 750 C. The water vapour addition was made by bubbling the gas through a water bath held at a controlled temperature of 46 C as illustrated in Fig. 2. The composition of the dry gas as guaranteed by the manufacturer is given in Table 2. The experiments for the dry atmosphere started after a flushing period of 12 h with synthetic air. The flow for the tests in moist air was opened during the heating stage at a temperature of 50 C. The samples were oxidized for 30, 100, 300 and 1000 h at 700 C and for 78 and 1000 h at 750 C. A switch of the atmosphere from dry to moist and back to dry conditions was conducted during the oxidation test at 750 C. Table 1 Chemical composition of the alloy Ti 47Al 1.0Cr Si Element
Ti
Al
Cr
Si
O
N
C
H
Wt.% At.%
64.25 51.42
33.25 47.24
1.45 1.07
1.12 0.16
0.02 0.047
0.005 0.013
0.008 0.025
0.002 0.075
Fig. 1. Geometry of as received ingots (broken line) and locations of the creep specimen and oxidation coupon (solid line).
61
A. Zeller et al. / Intermetallics 10 (2002) 59–72
Fig. 2. Experimental setup for the thermogravimetric measurements under dry and moist atmospheres.
X-ray diffraction analysis (XRD) of the oxidized surfaces was performed for phase identification. Metallographic cross-sections of all samples were prepared and investigated analytically using scanning electron microscopy (SEM) and electron probe microanalysis (EPMA). Prior to embedding the samples a nickel layer was galvanically deposited on the surface. The metal/scale interface of one sample (1000 h, 10% H2O, 700 C) was studied in greater detail using cross-sectional and energy-filtered transmission electron microscopy (XTEM/EFTEM). In addition, the oxide scales developed during the creep and fatigue tests were characterized by metallographic cross sections and SEM.
Table 2 Composition of the synthetic air N2
O2
Water vapour
CO2
NOx
Hydrocarbons
79.5 vol.%
20.5 vol.%
<5 vpm
<0.5 vpm
<0.1 vpm
<0.1 vpm
3. Results 3.1. Oxidation at 700 C 3.1.1. Kinetics The normalized mass gain as a function of oxidation time is shown in Fig. 3 for the two atmospheres. The mass gains discontinuously measured for the samples with shorter oxidation times (30, 100, and 300 h) fitted well to the 1000 h curves indicating a good reproducibility of the data. The oxidation kinetics was nearly linear for the
Fig. 3. Normalized mass gain as function of oxidation time for dry air and air with 10 vol.% H2O, 700 C showing accelerated and linear oxidation rates.
62
A. Zeller et al. / Intermetallics 10 (2002) 59–72
dry air with a rate constant of kl ¼ 2:6107 mg cm2 s1. The mass gain for the moist atmosphere is larger from the very beginning of the test with an accelerated growth up to about 600 h from where on the kinetics became linear with kl ¼ 7:1107 mg cm2 s1. The normalized mass gain was higher under moist conditions during the whole exposure test. 3.1.2. Microstructure of scale and metal/oxide interface The microstructural development of the sample surface is shown in Fig. 4 for both atmospheres after 30, 100 and 300 h oxidations. The XRD analysis revealed in
both cases only TiO2 (rutile) and a-Al2O3 as oxide phases. The presence of water vapour led, however, to a finer crystal morphology in the shape of needles and platelets. This difference became more obvious after the longest exposure time of 1000 h (Fig. 5). Here in dry air the outer oxide scale consisted mainly of TiO2-crystals with a ‘‘blocky‘‘ morphology. In contrast to the dry air, fine oxide needles developed under the moist conditions. Locally, the needles joined and grew together forming plates. The energy dispersive analysis (EDX) in the SEM revealed a ratio of Ti/Al=2/1 for the surface developed under dry conditions and of Ti/Al=1/1.25
Fig. 4. Different crystal morphologies developed on sample surfaces oxidized for (a) 30 h, (b) 100 h, and (c) 300 h in air and air+10 vol.% H2O.
A. Zeller et al. / Intermetallics 10 (2002) 59–72
63
for the moist conditions, indicating an enrichment of alumina near the surface of the scale due to the presence of water vapour. For the shorter oxidation times up to 300 h no significant differences in the morphologies of the oxide scales were observed from the metallographic cross sections. The cross sections revealed a mixture of titania and alumina and the development of an Al-depleted layer beneath the oxide scale. However, after 1000 h oxidation differences of the oxide scales became evident as shown in Fig. 6. As already described, in dry air the outer part of the scale consisted mainly of blocky rutile crystals whereas in moist air very fine and needle shaped oxide crystals formed on the outer part. The distribution of the two oxide phases, alumina and titania, in the scale was also influenced by the water vapour. In dry air a layered structure of the scale was formed where the outer TiO2 layer was followed by an alumina rich layer and the inward growing part of the scale. The latter was also layered containing alumina and titania rich layers but was dominated by the rutile phase. No specific Cr segregation/enrichment was measured within the resolution limit of the EPMA. The scale formed in the moist atmosphere did not show the outer coarse grained rutile
layer but a mixture of titania and alumina. The inner part of the scale consisted of a 5–6 mm thick titania rich layer followed by a 2–3 mm thick alumina rich layer. Some alumina islands and pores were present in the titania rich inner layer. Another difference in the scale microstructure between the two atmospheres occurred at the metal/ scale interface where under moist conditions an additional layer was present after longer oxidation times. This layer could be clearly distinguished from the Al-depleted layer and the oxide scale and is marked as ‘‘oxygen rich’’ layer in Fig. 6. The thickness of the oxide scale and the Aldepleted layer, respectively, as a function of oxidation time is given in Fig. 7. The measurements do not include the oxygen-rich layer (1.8 mm) of Fig. 6b. This is one reason why in the diagram the values for the scale thickness after 1000 h are similar for both atmospheres. In addition, the outer fine needles (Fig. 5b) were no longer visible in the metallographic cross section and may be lost during preparation. Furthermore, the oxide scale formed in the water vapour containing atmosphere had fewer pores and therefore a higher density compared to the scale formed under dry conditions which makes a comparison solely based on the measured thickness values difficult.
Fig. 5. Different crystal morphology of outer rutile layer after 1000 h oxidation at 700 C, (a) dry air; (b) air with 10 vol.% H2O.
Fig. 6. Cross-sections of oxide scales after 1000 h at 700 C, (a) dry air; (b) air+10 vol.% H2O (BSE mode) showing different distribution of the two oxide phases and different structure of the subsurface zones.
64
A. Zeller et al. / Intermetallics 10 (2002) 59–72
The thickness of the depletion layer increased with oxidation times up to 300 h after which a constant value was reached. In dry air this value is about 1.0 mm and for moist air 1.3 mm. If one takes into account the oxygenrich layer, a value of 3.1 mm has to be taken for the oxygen affected subsurface zone after 1000 h. 3.1.3. TEM investigations of the subsurface zone A TEM cross section of the sample oxidized for 1000 h in moist air was prepared to get further insight into the nature of the subsurface zone, especially the layer labelled as oxygen rich in Fig. 6b. Fig. 8 shows the TEM bright field images of the oxygen-rich and Al depletion layers. The Al depletion layer has been identified, by selected area electron diffraction (SAD), as the cubic phase with a ¼ 0:69 nm known from previous investigations [12–17]. Above this layer a two phase layer was present with a lamellar structure (Fig. 8a). The chemical composition of this layer was further investigated using electron spectroscopic imaging (ESI) [18]. The corresponding elemental maps given in Fig. 9 show that the lamellae consist of alumina next to Ti(Al,O). The two phases could be identified as a-Al2O3 and a-Ti by SAD. The structure of this layer is similar to the microstructure of
Fig. 7. Average thickness of oxide scale (a) and Al-depleted layer (b) as a function of oxidation time and atmosphere, 700 C.
the subsurface zone of a2-Ti3Al oxidized between 800 and 1000 C as reported recently in [19]. 3.2. Oxidation at 750 C 3.2.1. Kinetics Fig. 10 shows the normalized mass gain as a function of time for the oxidation at 750 C. The kinetics in dry air could be separated essentially into two parts with linear growth kinetics. A fast oxidation stage with kl ¼ 4:3106 mg cm2 s1 is followed after about 200 h by slower kinetics with kl ¼ 2:2107 mg cm2 s1. The value is comparable to the kinetics in dry air at 700 C. After 700 h a slight acceleration in the oxidation rate could be observed. The kinetics for the oxidation in air with 10 vol.% H2O showed a very different behaviour with an acceleration of
Fig. 8. BF-TEM images of the subsurface zone of sample shown in Fig. 6b (700 C, moist air, 1000 h). The oxygen rich layer consisits of a-Ti+a-Al2O3 (a) and is followed by the cubic phase and the base metal (b).
A. Zeller et al. / Intermetallics 10 (2002) 59–72
65
Fig. 9. Higher magnification of Fig. 8a and corresponding elemental maps of Ti, Al, and O. The bright appearing lamellae correspond to Al2O3, the darker ones to Ti(O,Al).
the oxidation rate from the very beginning reaching a linear rate of kl ¼ 2:5106 mg cm2 s1 after 300 h. The normalized mass gain and oxidation rate in moist air is in the early stage lower compared to the dry atmosphere. After 500 h the normalized mass gains are similar for both atmospheres but the oxidation rates are very different. The influence of a switch from dry to moist and back to dry air is illustrated in Fig. 11. The data have been corrected with respect to changes of the buoyancy effect due to the different atmospheres. The first part of the oxidation kinetics up to 45 h is similar to Fig. 3 for the dry atmosphere. The change from dry air to air with 10 vol.% H2O after 45 h did not lead to an instantaneous change in the oxidation kinetics. Instead a significant acceleration in the oxidation kinetics was obtained after an incubation period of about 20 h with an even slightly reduced oxidation rate. The oxidation rate was linear after about 100 h with a value of kl ¼ 6:9106 cm2 s1 and higher compared to the rates obtained for dry or moist atmospheres alone. After switching back to the dry air after 235 h the oxidation rate was immediately reduced and mostly decreasing until the end of the experiment (500 h).
3.2.2. Microstructure of oxide scale and metal/oxide interface The different surface morphologies of the oxide scales developed after 78 h at 750 C are shown in Fig. 12. The outer scale consisted of blocky rutile crystals comparable to the morphology observed at 700 C after 1000 h
Fig. 10. Normalized mass gain as a function of oxidation time for dry air and air with 10 vol.% H2O, 750 C.
66
A. Zeller et al. / Intermetallics 10 (2002) 59–72
(Fig. 5a). In moist air the outer scale was inhomogeneous with a mixture of needle-shaped and blocky TiO2 crystals. The areas with the needle-shaped crystals were comparable to the surface obtained after 1000 h at 700 C in moist air (Fig. 5b). After 1000 h the surface of the oxide scale was covered with TiO2 crystals which were cuboids for the moist air whereas in dry air the rutile crystals were more irregular in shape (Fig. 13). The rutile crystals of the outer oxide scale formed in dry air (Fig. 13a) showed no specific orientation and were also comparable with the surface of the sample oxidized at 700 C in dry air (Fig. 5a). The oxidation in moist air led instead to a preferred growth direction and morphology of the outer rutile crystals which were oriented mainly perpendicular to the original surface (Fig. 13b). Therefore, the presence of water vapour seems to favour the growth of specific crystal planes and growth directions of the rutile crystals. The difference in the scale microstructure also became evident in the cross-sections of the samples oxidized for
Fig. 11. Influence of a change from dry to moist and back to dry air on the oxidation kinetics, 750 C. The switch to moist air resulted in an accelerated oxidation after an incubation time of about 20 h.
1000 h. Fig. 14 shows the backscattered electron images of the corresponding cross sections. In dry air a nearly closed Al2O3 barrier was present beneath the outer TiO2 layer followed by the fine grained mixture of titania and alumina and the Al depletion layer. The average thickness of the scale was about 45 mm. This microstructure is similar to oxide scales obtained at higher temperatures (800– 900 C) and shorter times [16,20,33]. However, in moist air a thicker (95 mm) and surprisingly more adherent scale with a different microstructure developed. The outer part of the scale also consisted of titania but with a preferred columnar structure. Beneath and probably also within this outer layer alumina was precipitated but without forming a pronounced barrier layer as obtained under dry conditions. Beneath this zone a nearly pure 10–15 mm thick titania layer was present followed by the fine grained porous mixture of alumina and titania (60 mm) and the ‘‘oxygen-rich‘‘ (3 mm) and Al-depleted (4 mm) subsurface layers. 3.2.3. Oxide scale developed during atmosphere switch The metallographic cross-section for the sample following the oxidation kinetics given in Fig. 11 revealed a microstructure as obtained under moist conditions but with differences in the dimensions of particular characteristic features. The oxide scale had an average thickness of 47 mm which is only half the thickness of the oxide scale formed after 1000 h in moist air although the normalized mass gains were nearly similar. This discrepancy can be explained by the high porosity observed for the scale formed in moist air which was not observed for the oxide scale formed during the atmosphere switch. The surface showed mainly two different crystal morphologies corresponding to structures observed in solely dry and moist conditions (Fig. 15). Fig. 16 shows the two microstructures of the outer layer which were found after the switch experiment on the same sample. Also here the blocky coarse-grained outer rutile layer is connected with alumina enrichment (precipitates) below, whereas on the areas with the outer rutile needles the
Fig. 12. Surface morphology after 78 h at 750 C, (a) dry air, (b) air with 10 vol.% H2O.
A. Zeller et al. / Intermetallics 10 (2002) 59–72
alumina phase can also be found within the outer rutile layer (Fig. 16b) without forming a pronounced Al2O3 barrier. 3.3. Oxide scale developed under mechanical loading The microstructure and development of the oxide scale and subsurface zone play an important role in the mechanical properties of the alloy. Therefore, the oxide scales formed on the creep and LCF specimens have been investigated as well. To exclude surface finish effects as observed for the oxidation of TiAl in oxygen [14] specimens with a similar CNC finish as the creep and fatigue specimens were oxidized without loading for comparison. However, the scale microstructure of the oxidized specimens with a CNC finish were similar to the ones obtained on the 500 grit SiC finished samples. A significant influence of the surface finish was not observed. Fig. 17 illustrates the influence of stress and atmosphere on the oxide scales obtained after the creep tests at 700 C. In dry air and creep stresses between 200 and 290 MPa a 1–2 mm thin Al2O3-layer formed on the metal. At the lower stress of 170 MPa nodules with
67
thicknesses up to 15 mm formed. The oxidation in moist air resulted in a similar dependence of the nodule formation on the creep stress. For stresses between 200 and 290 MPa the thickness of the oxide scale was not homogeneous but thin areas were locally present. Compared to the oxide scales formed in dry air the oxidation was generally stronger. In both atmospheres the formation of protective alumina was (locally) observed for stress values between 200 and 290 MPa. The thickness of the oxide scale and Al-depletion layer was measured from the metallographic cross-sections. Because of the inhomogeneous thickness of some scales the highest values measured are given. The thickness values were highest for the nodule formation (170 MPa, dry air and 230 MPa, moist air). The oxide thickness obtained in dry air was lowest in the medium stresses between 200 and 290 MPa whereas no influence of stress on the thickness of the depletion layer was observed.
4. Discussion 4.1. General remarks The experimental results show that the oxidation behaviour of the alloy investigated was different for dry air compared to moist air. The presence of water vapour led to a change in the oxidation kinetics, the surface
Fig. 13. Sample plan views after oxidation for 1000 h at 750 C, (a) dry air, (b) air+10 vol.% H2O. In moist air the growth of specific crystal planes is favoured.
Fig. 14. Cross-sections of oxide scales after 1000 h at 750 C, (a) dry air, (b) air+10 vol.% H2O. The distribution of the oxide phases as well as the subsurface zone are different (SEM).
68
A. Zeller et al. / Intermetallics 10 (2002) 59–72
morphology, and the microstructure of the oxide scales formed. For a discussion of the mechanism several aspects have to be considered. In general a reaction of the water molecule with either the metal surface and/or with the oxide scale formed at high temperatures can take place. In case of titanium aluminides for the metal surface a reaction with the Al as partner is proposed in the literature for room temperature [20,21]. At high temperatures the reaction with the oxide phase(s) formed and the resulting formation of H atoms and/or hydroxide have
to be considered [22,23]. Especially the kind and location of the reactions are important for the oxidation behaviour [22]. Kremer et al. [5,6] proposed a reaction of the oxide with the water vapour as the important reaction for the oxidation of TiAl at 900 C. Motte et al. [24] observed a change in the oxidation behaviour of pure Ti and Ti6Al4V due to a surface reaction of water vapour and TiO2. An increasing water vapour partial pressure resulted in an increase of the oxidation kinetics whereby the Al containing alloy developed slower oxidation
Fig. 15. Different crystal morphologies developed on specimen surface after the switch experiment, 45 h dry air+190 h moist air+265 h dry air.
69
A. Zeller et al. / Intermetallics 10 (2002) 59–72
kinetics. Increasing the Al content further and comparing the oxidation kinetics of a2-Ti3Al in dry and moist air with the results obtained by Motte for pure Ti also shows that the increase in the kinetics due to water vapour is less for higher Al contents of the alloy [25]. Finally, for TiAl alloys which formed a protective alumina scale no significant effect of water vapour on the oxidation kinetics was reported [5,8,9]. These results suggest that the influence of water vapour strongly depends on the volume contents of TiO2 and Al2O3 in the oxide scale. Rutile and a-Al2O3 were the only oxide phases identified in dry as well as in moist air which indicates that no formation of metastable phases or other modifications is triggered by the presence of water vapour. Since the diffusivity of hydrogen in a-Al2O3 is 4–5 orders of magnitude slower compared to the TiO2 (rutile) (Table 3) it can be assumed that transport of H atoms mainly occurs via the rutile phase. In the following the interaction of water vapour with the rutile phase will be discussed in more detail since this reaction seems to be the most important for the oxidation behaviour of TiAl. The different surface morphologies observed after 1000 h oxidation can be explained by a favored growth of specific crystallographic planes. This can be related to adsorption studies of H2O molecules on rutile who showed that the dissociation of H2O preferably occurs on the (110) planes via the formation of a free H atom and an OH group [26,27] as schematically illustrated in Fig. 18. The presence of lattice defects enhances the preferential adsorption on the (110) planes even more. After the dissociation step the H preferably diffuses in
Table 3 Diffusion coefficient for H in Al2O3 and TiO2 (rutile) at 600–700 C Oxide
Direction
D (cm2/s)
T ( C)
Ref.
Al2O3 TiO2 TiO2
All directions ? c-axis k c-axis
1012–1011 108–107 106–5105
600–700 600–700 600–700
[38] [39] [39]
the c-direction via the oxygen atoms along the channels of the rutile crystal structure [28,29]. 4.2. Influence of water vapour In both atmospheres only small differences in the oxidation behaviour occurred for short oxidation times up to 300 h, mainly reflected in accelerated kinetics for the moist atmosphere. The finer needle structure observed after 300 h in moist air compared to dry air can not explain the higher mass gains. Also the thickness and porosity of the two scales are similar so that the only possible explanation for the difference in the mass gains is that the scales differ in their density. This means that the volume contents of the two oxide phases, rutile and alumina, have to be affected by the water vapour. Since the density of the rutile phase (TiO2=4.23–5.50 mg/cm3 [30]) is higher than that of the a-Al2O3 phase (Al2O3= 3.98–4.02 mg/cm3 [30]) the rutile formation and/or growth in moist air is enhanced. Together with the formation of the finer crystals the continuous increase in the oxidation kinetics during the oxidation in moist air could be caused by the faster growth of the rutile crystals and
Fig. 16. Regions with different oxide phase distributions within the outer scale obtained after the switch experiment, 45 h dry air+190 h moist air+265 h dry air.
70
A. Zeller et al. / Intermetallics 10 (2002) 59–72
the increase in the surface area for the reaction of the water molecules. This increase in the kinetics was not observed in dry air. However, a quantitative evaluation of the volume contents of the two oxide phases was not possible owing to the small dimensions of the scales after the shorter oxidation times. After 1000 h, pronounced differences in the surface morphology, oxide scale as well as subsurface zone occurred. The difference in the microstructure of the oxide scales indicates that the dissolution of H atoms or OHgroups in the rutile lattice led to a change in the defect chemistry and therefore in the transport and growth processes of the rutile phase. The preferred directional growth combined with the favoured growth of specific rutile crystal planes can be triggered by different adsorption kinetics for the H2O molecule on the various crystal planes and/or a preferred diffusion of the dissolved H/
OH species and, connected with it, of Ti and O ions in the rutile lattice. The difference in the alumina distribution/precipitation in the oxide scales could be explained by a change in the solubility of alumina in the rutile phase due to the presence of water vapour. Although the growth mechanism of the rather complex microstructure of the typical oxide scale obtained during the oxidation of TiAl in air is still not fully understood, one important mechanism could be the dissolution and reprecipitation of alumina in the rutile phase as dicussed by several authors [6,31–36]. Particularly concerning the influence of water vapour Sengupta et al. [37] found that a change in the defect structure of Al-doped rutile from n- to ptype is caused by adsorption of water vapour. This could explain why the alumina precipitates in the outer rutile-dominated layer in moist air whereas in dry air
Fig. 17. Oxide scales after creep tests. Stress axis was in horizontal direction (surfaces in SEM and metallographic cross-sections).
A. Zeller et al. / Intermetallics 10 (2002) 59–72
Fig. 18. Schematic diagram of the two-step adsorption process of H2O on the rutile (110) surface [26,27].
the precipitation occurs below the outer rutile layer. In addition, a change in the defect structure can also lead to an enhanced ion-diffusivity and therefore to accelerated oxidation kinetics. The defect structure of the rutile phase depends on the oxygen partial pressure and a change from p-type (high pO2) to n-type at lower pO2 occurs along the oxide scale. The solubility of Al interstitials may become less with increasing pO2 and precipitation of alumina occurs. If the defect structure is changed by the dissolution of H2O in the rutile phase a change in the precipitation behaviour of the alumina can be expected. However, more data of the rutile defect structure, the dissolution behaviour of H2O and Al2O3 and their influence on the rutile defect strucutre are needed to support this mechanism. Besides the outer layer which is mainly formed by outward cation diffusion and seems to be influenced by the water vapour reactions with the rutile, the metal/ scale interface is the second important location for the growth of the oxide scale since the inner part of the oxide scale grows by the inward transport of oxygen. Here a layer consisting of Al2O3 and a-Ti lamellae was formed in water vapour containing atmospheres. If the subsurface zone and the formation of this layer is influenced by reaction and/or dissolution of water vapour and/or hydrogen could not be proven. H-analysis performed after the oxidation experiments showed no significant increase of the H-content in the alloy [11]. However, the acting possible mechanisms could be a change in the stability of the cubic phase triggered by dissolution processes followed by its decomposition to Ti+Al2O3 or a change in the solubility/diffusivity of oxygen in the subsurface zone followed by internal oxidation. Another possibility
71
would be that the structure of the subsurface zone is only a result of the faster oxidation kinetics and Al depletion processes at the metal/scale interface and that the observed lamellae are the result of a ‘‘later oxidation’’ stage compared to the slower growing scale in dry air. Clarification of these points can not be obtained from the present results. The TGA data measured during the atmosphere changes from dry to moist and back to dry conditions showed marked changes of the oxidation kinetics. After an incubation time of about 20 h the oxidation kinetics was accelerated by changing the atmosphere from dry to moist air. Since the measured mass gain is mainly caused by the uptake of oxygen from the atmosphere the reaction of H2O during the incubation time must be limited to the surface of the scale because a significant additional uptake of water vapour immediately after the switch would lead to an acceleration of the kinetics compared to dry air from the time of the switch on. Therefore, it can be assumed that during the incubation period mainly a change in the defect structure of the rutile phase throughout the scale thickness occurs until the accelerated diffusion of ions is enabled. The switch back from moist to dry conditions led to an immediate change of the oxidation kinetics which can be explained by a stop of the H2O/TiO2 reaction and due to the high diffusivity of H by a fast change in the defect structure. Therefore, the slow down of the oxidation kinetics occurs without any significant incubation time. The creep stress and creep rate, respectively, had some influence on the thickness of the formed oxide scales during the creep tests. Especially at the lowest and highest stress levels in dry air thicker scales were locally formed on the creep specimens. This could be explained by formation of high growth stress for the low stress value and by significant scale cracking at the highest creep rate. For the medium stress values a balance between growth stress and stress release due to creep is obtained, favoring the formation of a thin alumina-based layer. However, as discussed in [11] the oxidation process was not the critical factor responsible for the accelerated creep in moist air.
5. Conclusions 1. The presence of water vapour changes in the atmosphere during the oxidation of TiAl led to a change in the oxidation kinetics, morphology and microstructure of the oxide scale as well as the subsurface zone. 2. The addition of water vapour led to an acceleration of the oxidation kinetics after a significant incubation time of about 20 h. Switching back from moist to dry air an immediate decrease in the kinetics was obtained.
72
A. Zeller et al. / Intermetallics 10 (2002) 59–72
3. Important parameters seem to be the reaction kinetics of water vapour with the rutile phase and a change in its defect chemistry leading to preferential crystal growth and morphologies of the rutile phase and to a different structure of the oxide scale. 4. The subsurface zone showed internal oxidation or the decomposition of the Al-depleted cubic phase for long time oxidation in moist atmosphere. If this is due to dissolution of H into the subsurface zone or due to the accelerated growth of the oxide scale and subsurface zone, respectively, could not be clarified yet. 5. As will be shown in part II of the work, the oxide scale and subsurface zone is not the critical feature concerning creep but play an important role for the fatigue properties of a component.
Acknowledgements The work presented was financially supported by the German Ministry of Economy and Technology (BMWi) via Arbeitsgemeinschaft industrieller Forschungsvereinigungen (AiF) under contract No. AiF 10898N, which is gratefully acknowledged. The authors would also like to thank Dr. W. Nisch, NMI Reutlingen for giving the opportunity to use the prototype of the new SESAMTEM at the University of Tu¨bingen developed by LEO Electron Microscopy Ltd and Mr. Nicolai, Tital GmbH for the supply of the material. Thanks are also due to Ms. Berghof-Hasselba¨cher (metallography), Ms. Schorr (EPMA), Mr. Gawenda (SEM), and Dr. Renusch (correction of text). References [1] Rahmel A, Tobolski J. Corros Sci 1965;5:333. [2] McCarron RL, Schulz JW. Proc. symp. on high temperature gas-metal reactions in mixed environments. New York: AIME, 1973. [3] Kofstad P. In: Bennett MJ, Lorimer GW, editors. Microscopy of oxidation. The Institute of Metals, London, 1991. p. 2. [4] Janakiraman R, Meier GH, Pettit FS. Metall Mater Trans A 1999;30A:2905. [5] Kremer R. Doctoral thesis, University of Erlangen-Nu¨rnberg, 1996. [6] Kremer R, Auer W. Mater Corros 1997;48:35. [7] Brady MP, Smialek JL, Humphrey DL, Smith J. Acta Mater 1997;45:2357.
[8] Brady MP, Smialek JL, Humphrey DL, Smith J. Acta Mater 1997;45:2371. [9] Hald M. Doctoral thesis, RWTH Aachen, Aachen, 1998. [10] Takasaki A, Ojima K, Taneda Y. Met Mater Trans 1994; A25:2491. [11] Zeller A, Dettenwanger F, Schu¨tze M. Part II of this study. Intermetallics. [12] Dowling WE, Donlon WT. Scripta Metall Mater 1992;27:1663. [13] Zheng N, Fischer W, Gru¨bmeier H, Shemet V, Quadakkers WJ. Scripta Metall Mater 1995;33:47. [14] Dettenwanger F, Schumann E, Ru¨hle M, Rakowski J, Meier GH. MRS Symp Proc 1995;364:981. [15] Cheng Y, Dettenwanger F, Mayer J, Schumann E, Ru¨hle M. Scripta Metall Mater 1996;34:707. [16] Lang C, Schu¨tze M. Oxid Met 1996;46:255. [17] Beye R, Verwerft M, Hosson JTMD, Gronsky R. Acta Mater 1996;44:4225. [18] Reimer L. In: Reimer L, editor. Energy-filtering transmission electron microscopy, Springer series in optical sciences 71. Springer, Berlin, Heidelberg, 1995. p. 347. [19] Dettenwanger F, Schu¨tze M. Oxid of Met 2000;54:121. [20] Stoloff NS, Shea M, Castagna A. In: Jones RH, Ricker RE, editors. Evironmental effects on advanced materials. The Minerals, Metals and Materials Society, 1991. p. 3. [21] Liu T. In: Izumi O, editor. Proc. of int. symp. on intermetallic compounds, structural and mechanical properties, JIMIS 6, The Japan Institute of Metals, 1991. p. 133. [22] Douglass DL, Kofstad P, Rahmel A, Wood GC. Oxid Met 1996; 45(5/6):529. [23] Kofstad P. High temperature corrosion. Essex, UK: Elsevier Applied Science, 1988. [24] Motte F, Coddet C, Sarrazin P, Azzopardi M, Besson J. Oxid Met 1976;10:113. [25] Dettenwanger F. Unpublished results. [26] Bredow T, Jug K. Surface Science 1995;327:398. [27] Henderson MA. Surface Science 1996;355:151. [28] Johnson OW, Paek SH, DeFord JW. J Appl Phys 1975; 46:1026. [29] Bates JB, Wang JC, Perkins RA. Phys Rev B 1979;19:4130. [30] Deer WA, Howie RA, Zussman J. The rock forming minerals. Essex, UK: Longman Scientific & Technical, 1992. [31] Kurtz RL, Stockbauer R, Madey TE, Roma´n E, DeSegovia JL. Surface Science 1989;218:178. [32] Slepetys RA, Vaughan PA. J Phys Chem 1969;73:2157. [33] Becker S, Rahmel A, Schorr M, Schu¨tze M. Oxid Met 1992; 38:425. [34] Lang C. Doctoral thesis, RWTH Aachen, Aachen, 1997. [35] Fergus JW, Dang B, Gale WF. Proc. 199th EFC meeting, 2001, Washington, DC, to be published. [36] Schumacher G, Dettenwanger F, Schu¨tze, Iberl A, Reil D. Oxid Met 2000;54:317. [37] Sengupta G, Chatterjee RN, Maity GC, Ansari BJ, Satyanarayna CVV. J Coll and Interf Sci 1995;170:215. [38] Fowler JD, Chandra D, Elleman TS, Payne AW, Verghese K. J Am Ceram Soc 1977;60:155. [39] Johnson OW, Paek SH, DeFord JW. J Appl Phys 1975; 46:1026.
Influence of water vapour on the oxidation behaviour of titanium aluminides A. Zeller, F. Dettenwanger, M. Schu¨tze* Karl-Winnacker-Institut der DECHEMA e.V., Theodor-Heuss-Allee 25, D-60486 Frankfurt am Main, Germany Received 29 May 2001; accepted 15 August 2001
Abstract The influence of water vapour on the oxidation of a technical g-TiAl based alloy at temperatures of 700 and 750 C is discussed. The present paper focuses on the oxidation behaviour in dry air and air containing 10 vol.% H2O with respect to oxidation kinetics and microstructural investigations of the oxide scales and the subsurface zones. The presence of water vapour led to a significant acceleration in the oxidation kinetics as well as to a change in the crystal morphology and microstructure of the oxide scales formed. The results indicate that the kinetics of the reaction of water molecules with the rutile phase and a change in its defect structure are important parameters in the mechanism. The subsurface zone formed in water vapour-containing atmospheres consisted (in addition to the cubic Al-depleted layer) of a lamellar layer of Al2O3 and Ti due to either internal oxidation or decomposition of the cubic phase. Both, the oxide scale and the depletion layer are important regarding the mechanical behaviour of a TiAl component at high temperatures which was also investigated. # 2002 Elsevier Science Ltd. All rights reserved. Keywords: A. Titanium aluminides based on TiAl; B. Oxidation; B. Phase identification; B. Environmental embrittlement
1. Introduction Titanium aluminides based on the intermetallic phase g-TiAl are attractive as a new group of high temperature alloys due to their specific high temperature strength. Research undertaken in the past 10 years was mainly concentrated on the understanding and improvement of the mechanical properties and oxidation behaviour of these alloys. However, the influence of water vapour on the oxidation behaviour and the mechanical properties of titanium aluminides has not been studied in the literature to a great extent yet although in the envisaged applications as gas turbine blades or car engine valves the atmosphere will contain a significant amount of H2O (up to 20 vol.%). In a more general sense water vapour can affect several properties of a high temperature component through the influence on the formation, microstructure, mechanical properties, adhesion, and transport properties of oxide scales [1–4] as well as the mechanical properties of the base metal directly. However, it has to be kept in mind that the influence of water vapour on the high * Corresponding author. Tel.: +49-69-7564-361; fax: +49-69-7564368.
temperature properties of alloys strongly depends on the specific system being considered, which makes a general modelling difficult if not impossible. A systematic study concerning the influence of water vapour on the oxidation of g-TiAl based alloys was conducted by Kremer et al. at 900 C in pure oxygen and oxygen with various water vapour contents [5,6]. The presence of water vapour led to a significant increase in the oxidation rate compared to the oxidation in dry oxygen. Interestingly, the accelerated kinetics was connected with an incubation time and the mass increase during the early stage was less in water vapour-containing atmosphere compared to dry conditions. After longer times ( 50 h) a linear oxidation rate was observed which was slowed down by switching the atmosphere to dry conditions. Changing back to the water vapourcontaining atmosphere accelerated the oxidation rate again after an incubation period. Variation of the water vapour and oxygen partial pressures revealed a variation of the linear rate constant with pH2O1/2 and pO21/4 which was interpreted with respect to different dissociation kinetics of the water molecule compared to the oxygen molecule. The different kinetics were reflected in the microstructures of the oxide scales formed. In dry oxygen
0966-9795/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(01)00104-2
60
A. Zeller et al. / Intermetallics 10 (2002) 59–72
a dense protective alumina barrier was formed whereas in humid oxygen only islands of alumina were present in the outer TiO2 layer. A possible mechanism of doping the oxide by hydrogen produced due to the dissociation of H2O was proposed [6]. Brady et al. [7,8] compared the oxidation behaviour of several TiAl,Cr-alloys with focus on the Cr-content at 1000 C in oxygen, and dry and wet air. In oxygen a dense alumina scale was observed on all alloys. In air the alloys with high Cr contents (15 at.%) consisting of g-TiAl and the Laves phase Ti(Cr,Al)2 formed a dense alumina scale independent of either dry or humidified air. This formation was promoted by the presence of the Laves phase. Therefore, an influence of water vapour was not observed for these alloys. In contrast, g-TiAl alloys with low Cr contents (< 5 at.%) showed a strong acceleration of the oxidation kinetics for the moist air. In addition, the structure of the subsurface zone could be correlated with the exposure conditions. The formation of a b-Ti(Al,Cr)-phase at the metal/oxide phase boundary hindering the formation of a protective alumina layer was observed for the moist atmosphere. It was concluded that the composition of the phase(s) present at the metal/scale interface is directly related to the influence of the water vapour in the atmosphere on the oxidation kinetics, i.e., on the formation of a protective alumina scale. Hald [9] observed an acceleration of the oxidation kinetics of TiAl at 900 C by switching from dry to moist air for the case of a non protective oxide scale on TiAl. For a previously formed alumina scale the switch to moist air did not lead to accelerated oxidation kinetics. Takasaki et al. [10] studied the influence of hydrogen dissolution in two phase titanium aluminides on the oxidation behaviour at 897 and 1077 C. The time of hydrogen loading influenced the subsequent oxidation behaviour significantly. Short loading times and therefore small amounts of dissolved hydrogen led to smaller weight gains compared to unloaded TiAl. Longer loading times led to higher weight gains which surprisingly increased with increasing Al content (Ti42Al, Ti45Al, Ti50Al). The beneficial effect of small hydrogen contents was attributed to a reduction of the Ti diffusion in the alloy, an explanation for the dependence of the Al content after long hydrogen loading times was not given. Investigations concerning the influence of water vapour on the mechanical and oxidation properties of gTiAl based alloys at the potential application temperatures of 700–750 C are to the authors’ knowledge not present in the literature. We investigated these points for a technical g-TiAl based alloy at temperatures of 700 and 750 C and the results will be presented in two papers. The present paper concentrates on the influence of water vapour on the oxidation behaviour and oxide scale microstructure for scales developed on the alloy with and without mechanical loading.
A second paper concentrates on the influence of water vapour on the creep and fatigue properties of the alloy Ti 47Al 1Cr Si and has been submitted to this journal.
2. Experimental The alloy Ti 47Al 1Cr Si investigated was provided by Tital, Bestwig, Germany, in the form of near specimen shaped slugs for the mechanical testing specimens in the as cast condition. The chemical composition is given in Table 1 and is close to the composition of a technical gbased TiAl alloy developed for application in automotive engines. The microstructure of the alloy is described in [11]. The material was cast in the near net shape of the cylindrical creep specimens. Fig. 1 illustrates geometry and locations of the oxidation and creep specimens cut from the as-received ingot. The oxidation specimens were cut from the lower part of the ingot with the dimensions 20102 mm3. The final grinding step for the coupons was performed with a 500 grit SiC paper and after grinding the coupons were ultrasonically cleaned in ethanol and acetone prior to the oxidation tests. The oxidation coupons were used for thermogravimetric measurements carried out in synthetic air and synthetic air containing 10 vol.% H2O at 700 and 750 C. The water vapour addition was made by bubbling the gas through a water bath held at a controlled temperature of 46 C as illustrated in Fig. 2. The composition of the dry gas as guaranteed by the manufacturer is given in Table 2. The experiments for the dry atmosphere started after a flushing period of 12 h with synthetic air. The flow for the tests in moist air was opened during the heating stage at a temperature of 50 C. The samples were oxidized for 30, 100, 300 and 1000 h at 700 C and for 78 and 1000 h at 750 C. A switch of the atmosphere from dry to moist and back to dry conditions was conducted during the oxidation test at 750 C. Table 1 Chemical composition of the alloy Ti 47Al 1.0Cr Si Element
Ti
Al
Cr
Si
O
N
C
H
Wt.% At.%
64.25 51.42
33.25 47.24
1.45 1.07
1.12 0.16
0.02 0.047
0.005 0.013
0.008 0.025
0.002 0.075
Fig. 1. Geometry of as received ingots (broken line) and locations of the creep specimen and oxidation coupon (solid line).
61
A. Zeller et al. / Intermetallics 10 (2002) 59–72
Fig. 2. Experimental setup for the thermogravimetric measurements under dry and moist atmospheres.
X-ray diffraction analysis (XRD) of the oxidized surfaces was performed for phase identification. Metallographic cross-sections of all samples were prepared and investigated analytically using scanning electron microscopy (SEM) and electron probe microanalysis (EPMA). Prior to embedding the samples a nickel layer was galvanically deposited on the surface. The metal/scale interface of one sample (1000 h, 10% H2O, 700 C) was studied in greater detail using cross-sectional and energy-filtered transmission electron microscopy (XTEM/EFTEM). In addition, the oxide scales developed during the creep and fatigue tests were characterized by metallographic cross sections and SEM.
Table 2 Composition of the synthetic air N2
O2
Water vapour
CO2
NOx
Hydrocarbons
79.5 vol.%
20.5 vol.%
<5 vpm
<0.5 vpm
<0.1 vpm
<0.1 vpm
3. Results 3.1. Oxidation at 700 C 3.1.1. Kinetics The normalized mass gain as a function of oxidation time is shown in Fig. 3 for the two atmospheres. The mass gains discontinuously measured for the samples with shorter oxidation times (30, 100, and 300 h) fitted well to the 1000 h curves indicating a good reproducibility of the data. The oxidation kinetics was nearly linear for the
Fig. 3. Normalized mass gain as function of oxidation time for dry air and air with 10 vol.% H2O, 700 C showing accelerated and linear oxidation rates.
62
A. Zeller et al. / Intermetallics 10 (2002) 59–72
dry air with a rate constant of kl ¼ 2:6107 mg cm2 s1. The mass gain for the moist atmosphere is larger from the very beginning of the test with an accelerated growth up to about 600 h from where on the kinetics became linear with kl ¼ 7:1107 mg cm2 s1. The normalized mass gain was higher under moist conditions during the whole exposure test. 3.1.2. Microstructure of scale and metal/oxide interface The microstructural development of the sample surface is shown in Fig. 4 for both atmospheres after 30, 100 and 300 h oxidations. The XRD analysis revealed in
both cases only TiO2 (rutile) and a-Al2O3 as oxide phases. The presence of water vapour led, however, to a finer crystal morphology in the shape of needles and platelets. This difference became more obvious after the longest exposure time of 1000 h (Fig. 5). Here in dry air the outer oxide scale consisted mainly of TiO2-crystals with a ‘‘blocky‘‘ morphology. In contrast to the dry air, fine oxide needles developed under the moist conditions. Locally, the needles joined and grew together forming plates. The energy dispersive analysis (EDX) in the SEM revealed a ratio of Ti/Al=2/1 for the surface developed under dry conditions and of Ti/Al=1/1.25
Fig. 4. Different crystal morphologies developed on sample surfaces oxidized for (a) 30 h, (b) 100 h, and (c) 300 h in air and air+10 vol.% H2O.
A. Zeller et al. / Intermetallics 10 (2002) 59–72
63
for the moist conditions, indicating an enrichment of alumina near the surface of the scale due to the presence of water vapour. For the shorter oxidation times up to 300 h no significant differences in the morphologies of the oxide scales were observed from the metallographic cross sections. The cross sections revealed a mixture of titania and alumina and the development of an Al-depleted layer beneath the oxide scale. However, after 1000 h oxidation differences of the oxide scales became evident as shown in Fig. 6. As already described, in dry air the outer part of the scale consisted mainly of blocky rutile crystals whereas in moist air very fine and needle shaped oxide crystals formed on the outer part. The distribution of the two oxide phases, alumina and titania, in the scale was also influenced by the water vapour. In dry air a layered structure of the scale was formed where the outer TiO2 layer was followed by an alumina rich layer and the inward growing part of the scale. The latter was also layered containing alumina and titania rich layers but was dominated by the rutile phase. No specific Cr segregation/enrichment was measured within the resolution limit of the EPMA. The scale formed in the moist atmosphere did not show the outer coarse grained rutile
layer but a mixture of titania and alumina. The inner part of the scale consisted of a 5–6 mm thick titania rich layer followed by a 2–3 mm thick alumina rich layer. Some alumina islands and pores were present in the titania rich inner layer. Another difference in the scale microstructure between the two atmospheres occurred at the metal/ scale interface where under moist conditions an additional layer was present after longer oxidation times. This layer could be clearly distinguished from the Al-depleted layer and the oxide scale and is marked as ‘‘oxygen rich’’ layer in Fig. 6. The thickness of the oxide scale and the Aldepleted layer, respectively, as a function of oxidation time is given in Fig. 7. The measurements do not include the oxygen-rich layer (1.8 mm) of Fig. 6b. This is one reason why in the diagram the values for the scale thickness after 1000 h are similar for both atmospheres. In addition, the outer fine needles (Fig. 5b) were no longer visible in the metallographic cross section and may be lost during preparation. Furthermore, the oxide scale formed in the water vapour containing atmosphere had fewer pores and therefore a higher density compared to the scale formed under dry conditions which makes a comparison solely based on the measured thickness values difficult.
Fig. 5. Different crystal morphology of outer rutile layer after 1000 h oxidation at 700 C, (a) dry air; (b) air with 10 vol.% H2O.
Fig. 6. Cross-sections of oxide scales after 1000 h at 700 C, (a) dry air; (b) air+10 vol.% H2O (BSE mode) showing different distribution of the two oxide phases and different structure of the subsurface zones.
64
A. Zeller et al. / Intermetallics 10 (2002) 59–72
The thickness of the depletion layer increased with oxidation times up to 300 h after which a constant value was reached. In dry air this value is about 1.0 mm and for moist air 1.3 mm. If one takes into account the oxygenrich layer, a value of 3.1 mm has to be taken for the oxygen affected subsurface zone after 1000 h. 3.1.3. TEM investigations of the subsurface zone A TEM cross section of the sample oxidized for 1000 h in moist air was prepared to get further insight into the nature of the subsurface zone, especially the layer labelled as oxygen rich in Fig. 6b. Fig. 8 shows the TEM bright field images of the oxygen-rich and Al depletion layers. The Al depletion layer has been identified, by selected area electron diffraction (SAD), as the cubic phase with a ¼ 0:69 nm known from previous investigations [12–17]. Above this layer a two phase layer was present with a lamellar structure (Fig. 8a). The chemical composition of this layer was further investigated using electron spectroscopic imaging (ESI) [18]. The corresponding elemental maps given in Fig. 9 show that the lamellae consist of alumina next to Ti(Al,O). The two phases could be identified as a-Al2O3 and a-Ti by SAD. The structure of this layer is similar to the microstructure of
Fig. 7. Average thickness of oxide scale (a) and Al-depleted layer (b) as a function of oxidation time and atmosphere, 700 C.
the subsurface zone of a2-Ti3Al oxidized between 800 and 1000 C as reported recently in [19]. 3.2. Oxidation at 750 C 3.2.1. Kinetics Fig. 10 shows the normalized mass gain as a function of time for the oxidation at 750 C. The kinetics in dry air could be separated essentially into two parts with linear growth kinetics. A fast oxidation stage with kl ¼ 4:3106 mg cm2 s1 is followed after about 200 h by slower kinetics with kl ¼ 2:2107 mg cm2 s1. The value is comparable to the kinetics in dry air at 700 C. After 700 h a slight acceleration in the oxidation rate could be observed. The kinetics for the oxidation in air with 10 vol.% H2O showed a very different behaviour with an acceleration of
Fig. 8. BF-TEM images of the subsurface zone of sample shown in Fig. 6b (700 C, moist air, 1000 h). The oxygen rich layer consisits of a-Ti+a-Al2O3 (a) and is followed by the cubic phase and the base metal (b).
A. Zeller et al. / Intermetallics 10 (2002) 59–72
65
Fig. 9. Higher magnification of Fig. 8a and corresponding elemental maps of Ti, Al, and O. The bright appearing lamellae correspond to Al2O3, the darker ones to Ti(O,Al).
the oxidation rate from the very beginning reaching a linear rate of kl ¼ 2:5106 mg cm2 s1 after 300 h. The normalized mass gain and oxidation rate in moist air is in the early stage lower compared to the dry atmosphere. After 500 h the normalized mass gains are similar for both atmospheres but the oxidation rates are very different. The influence of a switch from dry to moist and back to dry air is illustrated in Fig. 11. The data have been corrected with respect to changes of the buoyancy effect due to the different atmospheres. The first part of the oxidation kinetics up to 45 h is similar to Fig. 3 for the dry atmosphere. The change from dry air to air with 10 vol.% H2O after 45 h did not lead to an instantaneous change in the oxidation kinetics. Instead a significant acceleration in the oxidation kinetics was obtained after an incubation period of about 20 h with an even slightly reduced oxidation rate. The oxidation rate was linear after about 100 h with a value of kl ¼ 6:9106 cm2 s1 and higher compared to the rates obtained for dry or moist atmospheres alone. After switching back to the dry air after 235 h the oxidation rate was immediately reduced and mostly decreasing until the end of the experiment (500 h).
3.2.2. Microstructure of oxide scale and metal/oxide interface The different surface morphologies of the oxide scales developed after 78 h at 750 C are shown in Fig. 12. The outer scale consisted of blocky rutile crystals comparable to the morphology observed at 700 C after 1000 h
Fig. 10. Normalized mass gain as a function of oxidation time for dry air and air with 10 vol.% H2O, 750 C.
66
A. Zeller et al. / Intermetallics 10 (2002) 59–72
(Fig. 5a). In moist air the outer scale was inhomogeneous with a mixture of needle-shaped and blocky TiO2 crystals. The areas with the needle-shaped crystals were comparable to the surface obtained after 1000 h at 700 C in moist air (Fig. 5b). After 1000 h the surface of the oxide scale was covered with TiO2 crystals which were cuboids for the moist air whereas in dry air the rutile crystals were more irregular in shape (Fig. 13). The rutile crystals of the outer oxide scale formed in dry air (Fig. 13a) showed no specific orientation and were also comparable with the surface of the sample oxidized at 700 C in dry air (Fig. 5a). The oxidation in moist air led instead to a preferred growth direction and morphology of the outer rutile crystals which were oriented mainly perpendicular to the original surface (Fig. 13b). Therefore, the presence of water vapour seems to favour the growth of specific crystal planes and growth directions of the rutile crystals. The difference in the scale microstructure also became evident in the cross-sections of the samples oxidized for
Fig. 11. Influence of a change from dry to moist and back to dry air on the oxidation kinetics, 750 C. The switch to moist air resulted in an accelerated oxidation after an incubation time of about 20 h.
1000 h. Fig. 14 shows the backscattered electron images of the corresponding cross sections. In dry air a nearly closed Al2O3 barrier was present beneath the outer TiO2 layer followed by the fine grained mixture of titania and alumina and the Al depletion layer. The average thickness of the scale was about 45 mm. This microstructure is similar to oxide scales obtained at higher temperatures (800– 900 C) and shorter times [16,20,33]. However, in moist air a thicker (95 mm) and surprisingly more adherent scale with a different microstructure developed. The outer part of the scale also consisted of titania but with a preferred columnar structure. Beneath and probably also within this outer layer alumina was precipitated but without forming a pronounced barrier layer as obtained under dry conditions. Beneath this zone a nearly pure 10–15 mm thick titania layer was present followed by the fine grained porous mixture of alumina and titania (60 mm) and the ‘‘oxygen-rich‘‘ (3 mm) and Al-depleted (4 mm) subsurface layers. 3.2.3. Oxide scale developed during atmosphere switch The metallographic cross-section for the sample following the oxidation kinetics given in Fig. 11 revealed a microstructure as obtained under moist conditions but with differences in the dimensions of particular characteristic features. The oxide scale had an average thickness of 47 mm which is only half the thickness of the oxide scale formed after 1000 h in moist air although the normalized mass gains were nearly similar. This discrepancy can be explained by the high porosity observed for the scale formed in moist air which was not observed for the oxide scale formed during the atmosphere switch. The surface showed mainly two different crystal morphologies corresponding to structures observed in solely dry and moist conditions (Fig. 15). Fig. 16 shows the two microstructures of the outer layer which were found after the switch experiment on the same sample. Also here the blocky coarse-grained outer rutile layer is connected with alumina enrichment (precipitates) below, whereas on the areas with the outer rutile needles the
Fig. 12. Surface morphology after 78 h at 750 C, (a) dry air, (b) air with 10 vol.% H2O.
A. Zeller et al. / Intermetallics 10 (2002) 59–72
alumina phase can also be found within the outer rutile layer (Fig. 16b) without forming a pronounced Al2O3 barrier. 3.3. Oxide scale developed under mechanical loading The microstructure and development of the oxide scale and subsurface zone play an important role in the mechanical properties of the alloy. Therefore, the oxide scales formed on the creep and LCF specimens have been investigated as well. To exclude surface finish effects as observed for the oxidation of TiAl in oxygen [14] specimens with a similar CNC finish as the creep and fatigue specimens were oxidized without loading for comparison. However, the scale microstructure of the oxidized specimens with a CNC finish were similar to the ones obtained on the 500 grit SiC finished samples. A significant influence of the surface finish was not observed. Fig. 17 illustrates the influence of stress and atmosphere on the oxide scales obtained after the creep tests at 700 C. In dry air and creep stresses between 200 and 290 MPa a 1–2 mm thin Al2O3-layer formed on the metal. At the lower stress of 170 MPa nodules with
67
thicknesses up to 15 mm formed. The oxidation in moist air resulted in a similar dependence of the nodule formation on the creep stress. For stresses between 200 and 290 MPa the thickness of the oxide scale was not homogeneous but thin areas were locally present. Compared to the oxide scales formed in dry air the oxidation was generally stronger. In both atmospheres the formation of protective alumina was (locally) observed for stress values between 200 and 290 MPa. The thickness of the oxide scale and Al-depletion layer was measured from the metallographic cross-sections. Because of the inhomogeneous thickness of some scales the highest values measured are given. The thickness values were highest for the nodule formation (170 MPa, dry air and 230 MPa, moist air). The oxide thickness obtained in dry air was lowest in the medium stresses between 200 and 290 MPa whereas no influence of stress on the thickness of the depletion layer was observed.
4. Discussion 4.1. General remarks The experimental results show that the oxidation behaviour of the alloy investigated was different for dry air compared to moist air. The presence of water vapour led to a change in the oxidation kinetics, the surface
Fig. 13. Sample plan views after oxidation for 1000 h at 750 C, (a) dry air, (b) air+10 vol.% H2O. In moist air the growth of specific crystal planes is favoured.
Fig. 14. Cross-sections of oxide scales after 1000 h at 750 C, (a) dry air, (b) air+10 vol.% H2O. The distribution of the oxide phases as well as the subsurface zone are different (SEM).
68
A. Zeller et al. / Intermetallics 10 (2002) 59–72
morphology, and the microstructure of the oxide scales formed. For a discussion of the mechanism several aspects have to be considered. In general a reaction of the water molecule with either the metal surface and/or with the oxide scale formed at high temperatures can take place. In case of titanium aluminides for the metal surface a reaction with the Al as partner is proposed in the literature for room temperature [20,21]. At high temperatures the reaction with the oxide phase(s) formed and the resulting formation of H atoms and/or hydroxide have
to be considered [22,23]. Especially the kind and location of the reactions are important for the oxidation behaviour [22]. Kremer et al. [5,6] proposed a reaction of the oxide with the water vapour as the important reaction for the oxidation of TiAl at 900 C. Motte et al. [24] observed a change in the oxidation behaviour of pure Ti and Ti6Al4V due to a surface reaction of water vapour and TiO2. An increasing water vapour partial pressure resulted in an increase of the oxidation kinetics whereby the Al containing alloy developed slower oxidation
Fig. 15. Different crystal morphologies developed on specimen surface after the switch experiment, 45 h dry air+190 h moist air+265 h dry air.
69
A. Zeller et al. / Intermetallics 10 (2002) 59–72
kinetics. Increasing the Al content further and comparing the oxidation kinetics of a2-Ti3Al in dry and moist air with the results obtained by Motte for pure Ti also shows that the increase in the kinetics due to water vapour is less for higher Al contents of the alloy [25]. Finally, for TiAl alloys which formed a protective alumina scale no significant effect of water vapour on the oxidation kinetics was reported [5,8,9]. These results suggest that the influence of water vapour strongly depends on the volume contents of TiO2 and Al2O3 in the oxide scale. Rutile and a-Al2O3 were the only oxide phases identified in dry as well as in moist air which indicates that no formation of metastable phases or other modifications is triggered by the presence of water vapour. Since the diffusivity of hydrogen in a-Al2O3 is 4–5 orders of magnitude slower compared to the TiO2 (rutile) (Table 3) it can be assumed that transport of H atoms mainly occurs via the rutile phase. In the following the interaction of water vapour with the rutile phase will be discussed in more detail since this reaction seems to be the most important for the oxidation behaviour of TiAl. The different surface morphologies observed after 1000 h oxidation can be explained by a favored growth of specific crystallographic planes. This can be related to adsorption studies of H2O molecules on rutile who showed that the dissociation of H2O preferably occurs on the (110) planes via the formation of a free H atom and an OH group [26,27] as schematically illustrated in Fig. 18. The presence of lattice defects enhances the preferential adsorption on the (110) planes even more. After the dissociation step the H preferably diffuses in
Table 3 Diffusion coefficient for H in Al2O3 and TiO2 (rutile) at 600–700 C Oxide
Direction
D (cm2/s)
T ( C)
Ref.
Al2O3 TiO2 TiO2
All directions ? c-axis k c-axis
1012–1011 108–107 106–5105
600–700 600–700 600–700
[38] [39] [39]
the c-direction via the oxygen atoms along the channels of the rutile crystal structure [28,29]. 4.2. Influence of water vapour In both atmospheres only small differences in the oxidation behaviour occurred for short oxidation times up to 300 h, mainly reflected in accelerated kinetics for the moist atmosphere. The finer needle structure observed after 300 h in moist air compared to dry air can not explain the higher mass gains. Also the thickness and porosity of the two scales are similar so that the only possible explanation for the difference in the mass gains is that the scales differ in their density. This means that the volume contents of the two oxide phases, rutile and alumina, have to be affected by the water vapour. Since the density of the rutile phase (TiO2=4.23–5.50 mg/cm3 [30]) is higher than that of the a-Al2O3 phase (Al2O3= 3.98–4.02 mg/cm3 [30]) the rutile formation and/or growth in moist air is enhanced. Together with the formation of the finer crystals the continuous increase in the oxidation kinetics during the oxidation in moist air could be caused by the faster growth of the rutile crystals and
Fig. 16. Regions with different oxide phase distributions within the outer scale obtained after the switch experiment, 45 h dry air+190 h moist air+265 h dry air.
70
A. Zeller et al. / Intermetallics 10 (2002) 59–72
the increase in the surface area for the reaction of the water molecules. This increase in the kinetics was not observed in dry air. However, a quantitative evaluation of the volume contents of the two oxide phases was not possible owing to the small dimensions of the scales after the shorter oxidation times. After 1000 h, pronounced differences in the surface morphology, oxide scale as well as subsurface zone occurred. The difference in the microstructure of the oxide scales indicates that the dissolution of H atoms or OHgroups in the rutile lattice led to a change in the defect chemistry and therefore in the transport and growth processes of the rutile phase. The preferred directional growth combined with the favoured growth of specific rutile crystal planes can be triggered by different adsorption kinetics for the H2O molecule on the various crystal planes and/or a preferred diffusion of the dissolved H/
OH species and, connected with it, of Ti and O ions in the rutile lattice. The difference in the alumina distribution/precipitation in the oxide scales could be explained by a change in the solubility of alumina in the rutile phase due to the presence of water vapour. Although the growth mechanism of the rather complex microstructure of the typical oxide scale obtained during the oxidation of TiAl in air is still not fully understood, one important mechanism could be the dissolution and reprecipitation of alumina in the rutile phase as dicussed by several authors [6,31–36]. Particularly concerning the influence of water vapour Sengupta et al. [37] found that a change in the defect structure of Al-doped rutile from n- to ptype is caused by adsorption of water vapour. This could explain why the alumina precipitates in the outer rutile-dominated layer in moist air whereas in dry air
Fig. 17. Oxide scales after creep tests. Stress axis was in horizontal direction (surfaces in SEM and metallographic cross-sections).
A. Zeller et al. / Intermetallics 10 (2002) 59–72
Fig. 18. Schematic diagram of the two-step adsorption process of H2O on the rutile (110) surface [26,27].
the precipitation occurs below the outer rutile layer. In addition, a change in the defect structure can also lead to an enhanced ion-diffusivity and therefore to accelerated oxidation kinetics. The defect structure of the rutile phase depends on the oxygen partial pressure and a change from p-type (high pO2) to n-type at lower pO2 occurs along the oxide scale. The solubility of Al interstitials may become less with increasing pO2 and precipitation of alumina occurs. If the defect structure is changed by the dissolution of H2O in the rutile phase a change in the precipitation behaviour of the alumina can be expected. However, more data of the rutile defect structure, the dissolution behaviour of H2O and Al2O3 and their influence on the rutile defect strucutre are needed to support this mechanism. Besides the outer layer which is mainly formed by outward cation diffusion and seems to be influenced by the water vapour reactions with the rutile, the metal/ scale interface is the second important location for the growth of the oxide scale since the inner part of the oxide scale grows by the inward transport of oxygen. Here a layer consisting of Al2O3 and a-Ti lamellae was formed in water vapour containing atmospheres. If the subsurface zone and the formation of this layer is influenced by reaction and/or dissolution of water vapour and/or hydrogen could not be proven. H-analysis performed after the oxidation experiments showed no significant increase of the H-content in the alloy [11]. However, the acting possible mechanisms could be a change in the stability of the cubic phase triggered by dissolution processes followed by its decomposition to Ti+Al2O3 or a change in the solubility/diffusivity of oxygen in the subsurface zone followed by internal oxidation. Another possibility
71
would be that the structure of the subsurface zone is only a result of the faster oxidation kinetics and Al depletion processes at the metal/scale interface and that the observed lamellae are the result of a ‘‘later oxidation’’ stage compared to the slower growing scale in dry air. Clarification of these points can not be obtained from the present results. The TGA data measured during the atmosphere changes from dry to moist and back to dry conditions showed marked changes of the oxidation kinetics. After an incubation time of about 20 h the oxidation kinetics was accelerated by changing the atmosphere from dry to moist air. Since the measured mass gain is mainly caused by the uptake of oxygen from the atmosphere the reaction of H2O during the incubation time must be limited to the surface of the scale because a significant additional uptake of water vapour immediately after the switch would lead to an acceleration of the kinetics compared to dry air from the time of the switch on. Therefore, it can be assumed that during the incubation period mainly a change in the defect structure of the rutile phase throughout the scale thickness occurs until the accelerated diffusion of ions is enabled. The switch back from moist to dry conditions led to an immediate change of the oxidation kinetics which can be explained by a stop of the H2O/TiO2 reaction and due to the high diffusivity of H by a fast change in the defect structure. Therefore, the slow down of the oxidation kinetics occurs without any significant incubation time. The creep stress and creep rate, respectively, had some influence on the thickness of the formed oxide scales during the creep tests. Especially at the lowest and highest stress levels in dry air thicker scales were locally formed on the creep specimens. This could be explained by formation of high growth stress for the low stress value and by significant scale cracking at the highest creep rate. For the medium stress values a balance between growth stress and stress release due to creep is obtained, favoring the formation of a thin alumina-based layer. However, as discussed in [11] the oxidation process was not the critical factor responsible for the accelerated creep in moist air.
5. Conclusions 1. The presence of water vapour changes in the atmosphere during the oxidation of TiAl led to a change in the oxidation kinetics, morphology and microstructure of the oxide scale as well as the subsurface zone. 2. The addition of water vapour led to an acceleration of the oxidation kinetics after a significant incubation time of about 20 h. Switching back from moist to dry air an immediate decrease in the kinetics was obtained.
72
A. Zeller et al. / Intermetallics 10 (2002) 59–72
3. Important parameters seem to be the reaction kinetics of water vapour with the rutile phase and a change in its defect chemistry leading to preferential crystal growth and morphologies of the rutile phase and to a different structure of the oxide scale. 4. The subsurface zone showed internal oxidation or the decomposition of the Al-depleted cubic phase for long time oxidation in moist atmosphere. If this is due to dissolution of H into the subsurface zone or due to the accelerated growth of the oxide scale and subsurface zone, respectively, could not be clarified yet. 5. As will be shown in part II of the work, the oxide scale and subsurface zone is not the critical feature concerning creep but play an important role for the fatigue properties of a component.
Acknowledgements The work presented was financially supported by the German Ministry of Economy and Technology (BMWi) via Arbeitsgemeinschaft industrieller Forschungsvereinigungen (AiF) under contract No. AiF 10898N, which is gratefully acknowledged. The authors would also like to thank Dr. W. Nisch, NMI Reutlingen for giving the opportunity to use the prototype of the new SESAMTEM at the University of Tu¨bingen developed by LEO Electron Microscopy Ltd and Mr. Nicolai, Tital GmbH for the supply of the material. Thanks are also due to Ms. Berghof-Hasselba¨cher (metallography), Ms. Schorr (EPMA), Mr. Gawenda (SEM), and Dr. Renusch (correction of text). References [1] Rahmel A, Tobolski J. Corros Sci 1965;5:333. [2] McCarron RL, Schulz JW. Proc. symp. on high temperature gas-metal reactions in mixed environments. New York: AIME, 1973. [3] Kofstad P. In: Bennett MJ, Lorimer GW, editors. Microscopy of oxidation. The Institute of Metals, London, 1991. p. 2. [4] Janakiraman R, Meier GH, Pettit FS. Metall Mater Trans A 1999;30A:2905. [5] Kremer R. Doctoral thesis, University of Erlangen-Nu¨rnberg, 1996. [6] Kremer R, Auer W. Mater Corros 1997;48:35. [7] Brady MP, Smialek JL, Humphrey DL, Smith J. Acta Mater 1997;45:2357.
[8] Brady MP, Smialek JL, Humphrey DL, Smith J. Acta Mater 1997;45:2371. [9] Hald M. Doctoral thesis, RWTH Aachen, Aachen, 1998. [10] Takasaki A, Ojima K, Taneda Y. Met Mater Trans 1994; A25:2491. [11] Zeller A, Dettenwanger F, Schu¨tze M. Part II of this study. Intermetallics. [12] Dowling WE, Donlon WT. Scripta Metall Mater 1992;27:1663. [13] Zheng N, Fischer W, Gru¨bmeier H, Shemet V, Quadakkers WJ. Scripta Metall Mater 1995;33:47. [14] Dettenwanger F, Schumann E, Ru¨hle M, Rakowski J, Meier GH. MRS Symp Proc 1995;364:981. [15] Cheng Y, Dettenwanger F, Mayer J, Schumann E, Ru¨hle M. Scripta Metall Mater 1996;34:707. [16] Lang C, Schu¨tze M. Oxid Met 1996;46:255. [17] Beye R, Verwerft M, Hosson JTMD, Gronsky R. Acta Mater 1996;44:4225. [18] Reimer L. In: Reimer L, editor. Energy-filtering transmission electron microscopy, Springer series in optical sciences 71. Springer, Berlin, Heidelberg, 1995. p. 347. [19] Dettenwanger F, Schu¨tze M. Oxid of Met 2000;54:121. [20] Stoloff NS, Shea M, Castagna A. In: Jones RH, Ricker RE, editors. Evironmental effects on advanced materials. The Minerals, Metals and Materials Society, 1991. p. 3. [21] Liu T. In: Izumi O, editor. Proc. of int. symp. on intermetallic compounds, structural and mechanical properties, JIMIS 6, The Japan Institute of Metals, 1991. p. 133. [22] Douglass DL, Kofstad P, Rahmel A, Wood GC. Oxid Met 1996; 45(5/6):529. [23] Kofstad P. High temperature corrosion. Essex, UK: Elsevier Applied Science, 1988. [24] Motte F, Coddet C, Sarrazin P, Azzopardi M, Besson J. Oxid Met 1976;10:113. [25] Dettenwanger F. Unpublished results. [26] Bredow T, Jug K. Surface Science 1995;327:398. [27] Henderson MA. Surface Science 1996;355:151. [28] Johnson OW, Paek SH, DeFord JW. J Appl Phys 1975; 46:1026. [29] Bates JB, Wang JC, Perkins RA. Phys Rev B 1979;19:4130. [30] Deer WA, Howie RA, Zussman J. The rock forming minerals. Essex, UK: Longman Scientific & Technical, 1992. [31] Kurtz RL, Stockbauer R, Madey TE, Roma´n E, DeSegovia JL. Surface Science 1989;218:178. [32] Slepetys RA, Vaughan PA. J Phys Chem 1969;73:2157. [33] Becker S, Rahmel A, Schorr M, Schu¨tze M. Oxid Met 1992; 38:425. [34] Lang C. Doctoral thesis, RWTH Aachen, Aachen, 1997. [35] Fergus JW, Dang B, Gale WF. Proc. 199th EFC meeting, 2001, Washington, DC, to be published. [36] Schumacher G, Dettenwanger F, Schu¨tze, Iberl A, Reil D. Oxid Met 2000;54:317. [37] Sengupta G, Chatterjee RN, Maity GC, Ansari BJ, Satyanarayna CVV. J Coll and Interf Sci 1995;170:215. [38] Fowler JD, Chandra D, Elleman TS, Payne AW, Verghese K. J Am Ceram Soc 1977;60:155. [39] Johnson OW, Paek SH, DeFord JW. J Appl Phys 1975; 46:1026.