- Email: [email protected]
Influence of sol–gel derived Al2O3 film on the oxidation behavior of a Ti3Al based alloy
Materials Science and Engineering A 415 (2006) 177–183
Influence of sol–gel derived Al2O3 film on the oxidation behavior of a Ti3Al based alloy Ming Zhu a,b , Meishuan Li a,∗ , Yali Li a , Yanchun Zhou a a
High-Performance Ceramic Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of Chinese Academy of Sciences, Beijing 100039, China Received in revised form 6 September 2005; accepted 20 September 2005
Abstract Al2 O3 thin films were deposited on a Ti3 Al based alloy (Ti–24Al–14Nb–3V–0.5Mo–0.3Si) by sol–gel processing. Isothermal oxidation at temperatures of 900–1000 ◦ C and cyclic oxidation at 800–900 ◦ C were performed to test their effect on the oxidation behavior of the alloy. Results of the oxidation tests show that the oxidation parabolic rate constants of the alloy were reduced due to the applied thin film. This beneficial effect became weaker after longer oxidation time at 1000 ◦ C. TiO2 and Al2 O3 were the main phases formed on the alloy. The thin film could promote the growth of Al2 O3 , causing an increase of the Al2 O3 content in the composite oxides, sequentially decreased the oxidation rate. Nb/Al enriched as a layer in the alloy adjacent to the oxide/alloy interface in both the coated and uncoated alloy. The coated thin film decreased the thickness of the Nb/Al enrichment layer by reducing the scale growth rate. © 2005 Elsevier B.V. All rights reserved. Keywords: Ti3 Al based alloy; Sol–gel processing; Al2 O3 thin film; High-temperature oxidation
1. Introduction Ti3 Al based alloys are prospective high-temperature structural materials for aeroengine and automobile parts because of their excellent mechanical properties such as high specific strength and high elastic modulus at elevated temperatures [1–3]. However, poor oxidation resistance of these materials at high temperatures, which is attributed to the formation of mixed oxides containing the less protective TiO2 (rutile), and ␣-Al2 O3 , rather than a continuous protective ␣-Al2 O3 scale in oxidizing environments, limits their industrial applications. Rahmel and Spencer [4] have calculated the Al- and Ti-activities (aAl and aTi ) in Ti–Al intermetallic compounds. Their results show that aluminum concentration of over 50 at.% is needed for the growth of a single phase external Al2 O3 layer. Additions of some transition elements, such as Nb [5–8], Ta, Cr [9], Mo, W and Re can increase the activity ratio (aAl /aTi ) in these alloys, causing the protective alumina scale to be formed even when the aluminum concentration is less than 50 at.% [10]. However, some of the
∗
Corresponding author. Tel.: +86 24 23915913; fax: +86 24 23891320. E-mail address: [email protected] (M. Li).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.09.078
added elements may have negative effects on the mechanical properties of the alloys. This problem may be solved by depositing coatings on the alloys because the application of coatings may improve their oxidation resistance without degrading their mechanical properties. Many kinds of the coatings have been prepared in order to improve the oxidation resistance of Ti3 Al based alloys, such as diffusing coatings [11–15] and depositing coatings [16–18]. All of these coatings are meant to form a continuous protective Al2 O3 or Cr2 O3 layer on the surface during oxidation. Aluminide coatings are quite effective to improve the oxidation resistance of Ti3 Al based alloys, however, cracks generate easily in the coatings [15] due to large thermal stresses resulting from the thermal expansion mismatch between the coatings and the substrate alloys. Most of the deposited coatings usually have shortcomings of poor adhesion [17] and significant inter diffusion [18] between the coating and the substrate. Thus, new technologies or new coating systems need to be developed in order to further improve the oxidation resistance of Ti–Al based alloys. Sol–gel processing is widely used to prepare various functional thin films in which high-temperature oxidation resistant coatings are included [19,20], but its use as a protective coating on Ti–Al based alloys has not been reported. The generation of
178
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
cracks during thermal densification of the gel at around 500 ◦ C, and the lack of adhesion between these films and metals limit the use of this technique in preparing thick high-temperature oxidation resistant coatings, but cracking is seldom a problem for thin films. The aim of the present work is to study the effect of sol–gel derived Al2 O3 thin film on the selective oxidation of a Ti3 Al based alloy (Ti–24Al–14Nb–3V–0.5Mo–0.3Si, at.%), and to probe into the feasibility of using this sol–gel derived thin film as an oxidation resistant coating on this alloy.
reached the desired value. Cyclic oxidation was conducted at 800 ◦ C and 900 ◦ C. The specimens were kept in the furnace at 800 or 900 ◦ C for 1 h and then taken out to cool for 10 min in air. This process is defined as one cycle, and the cycle was repeated a few 100 times. An analytical balance with a sensitivity of 10−5 g was used to measure the weight of the specimens at intervals during the cyclic oxidation test. The surface morphologies and cross-sections of the oxidized samples were characterized by SEM equipped with EDS, while the phases of the oxide scales were analyzed using an X-ray diffractometer with Cu K␣ radiation.
2. Experimental 3. Results The nominal composition of the Ti3 Al based alloy is Ti–24Al–14Nb–3V–0.5Mo–0.3Si (at.%). The alloy ingot was prepared by melting twice in a consumable-electrode arc furnace under vacuum. Billets were cut from the ingots and, after descaling, hot press forged first in the -phase region at 1200 ◦ C and then in the (␣2 + 2 ) two-phase region at 1050 ◦ C with intermediate annealing. The heat treatment schedule of the alloy after forging was 1000 ◦ C/1 h/AC + 850 ◦ C/2 h/AC (AC stands for air-cooling). The alloy consists of three phases (␣2 + 2 + O). Specimens were cut into dimensions of 10 mm × 12 mm × 2 mm by using a spark wire machine, and their surfaces were grounded to 600 grit SiC paper. In order to improve the adhesion between the thin film and the substrate, the surfaces of the substrate were treated by sandblasting with steel abrasive powders. Then the samples were degreased ultrasonically in acetone, cleaned by distilled water and dried in air. Al2 O3 sol was prepared through the method used by Yoldas [21]. 20.4 g aluminum isopropoxide (Al(OC3 H7 )3 ), purity >98%, was hydrated in 270 ml distilled water at about 80 ◦ C with stirring for 1 h, then HNO3 was added into the solution to adjust the pH to about 3. The vessel containing the solution was kept constantly at 90 ◦ C for 16 h; a transparent sol was obtained. Samples were dipped into the sol and withdrawn at a rate between 3 and 30 cm/min, then dried in air for 15 min at room temperature. Thicker films could be obtained by repeating the dipping step. In this work samples were dipped and dried for five times. Thermal treatment was carried out in a tube furnace. The tube was first evacuated to a pressure of 2 × 10−1 Pa and then argon was introduced as a protective gas at a flow rate of 0.5 l/min. The tube was heated up to 500 ◦ C at the rate of 4 ◦ C/min and kept at this temperature for half an hour. Then it was heat up to 1000 ◦ C at the same heating rate and kept for 30 min. The target sol–gel Al2 O3 thin films were finally obtained by cooling the specimens to ambient temperature in the furnace. In order to understand the structure transformation of the gel during the annealing process, the dry gel was heated at 500, 700, 900 and 1000 ◦ C for half an hour, respectively, then grinded into powders and sieved. XRD powder analysis was used to examine the structure of the annealed gels. The continuous oxidation weight gain measurements were performed in a Sytses16/18(Setaram, France) microbalance at temperatures of 900–1000 ◦ C in air for 20 h. The samples were heated at a rate of 40 ◦ C/min to the test temperature. The mass change was recorded as a function of time once the temperature
3.1. Characteristics of the sol XRD patterns of the Al2 O3 gels treated for half an hour at 500–1000 ◦ C and then grinded into powers are presented in Fig. 1. Three main peaks of ␥-Al2 O3 were detected on the samples annealed at 500 and 700 ◦ C. The peaks corresponding to ␦-Al2 O3 also appeared after the treatment at 900 ◦ C. The gels treated below 1000 ◦ C was not well crystallized. The gel treated at 1000 ◦ C composed of ␣-Al2 O3 and a small mount of -Al2 O3 . Since a protective ␣-Al2 O3 layer is expected to provide the oxidation protection, the annealing temperature of the coating cannot be lower than 1000 ◦ C. But considering that the substrate alloy was heat treated at the temperature of 1000 ◦ C, the final annealing temperature of the coatings was better not to above 1000 ◦ C. 3.2. Oxidation kinetics The isothermal oxidation kinetics at different temperatures of the blank and coated samples are depicted in Fig. 2 as a function of the square of mass gain versus oxidation time. It can be seen that at the same oxidation temperature, the weight gains of the Al2 O3 coated samples were lower than those of Ti3 Al based alloy. The oxidation rate of the coated sample at 1000 ◦ C
Fig. 1. XRD patterns of the Al2 O3 gels after drying at (a) 500 ◦ C, (b) 700 ◦ C, (c) 900 ◦ C, and (d) 1000 ◦ C for 30 min.
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
Fig. 2. Plots of square of mass gain vs. time during the isothermal oxidation of blank and coated Ti3 Al based alloy at 900 and 1000 ◦ C. Table 1 Parabolic rate constants for different samples oxidized at 900 and 1000 ◦ C Parabolic rate constant, kp (mg2 cm−4 s−1 ) 900 ◦ C
1000 ◦ C
Blank alloy
2.20 × 10−4
8.32 × 10−4 (0–4 h) 3.96 × 10−4 (4–20 h)
Coated alloy
2.68 × 10−5
1.67 × 10−4
was even slower than that of the blank sample at 900 ◦ C. The oxidation kinetic of the blank sample at 1000 ◦ C deviated from the parabolic law while the other three curves presented in Fig. 2 followed the parabolic law after the first few hours. The parabolic rate constant (kp ) could be obtained by a linear fit of m2 = A + kp t, where A is a constant and the kp values are listed in Table 1. It is apparent that when coated with Al2 O3 film, the parabolic rate constants of the alloy decreased about 88% at 900 ◦ C, and roughly 60% at 1000 ◦ C. Fig. 3 shows the cyclic oxidation kinetics of the blank and coated samples at 800 ◦ C (a) and 900 ◦ C (b). The reason the cyclic oxidation was carried out at 800 and 900 ◦ C was because
179
the temperature at which Ti3 Al based alloys used in practical applications are usually between 800 and 900 ◦ C. It can be seen that the mass gain of each sample generally kept increasing during the cyclic oxidation. The oxidation rate of the coated sample was much lower than that of the blank sample when cyclic oxidized at 800 ◦ C (Fig. 3(a)). During the cyclic oxidation at 900 ◦ C (Fig. 3(b)), the sol–gel derived Al2 O3 thin film did not decrease the oxidation rate as remarkably as it did at 800 ◦ C. The curve corresponding to the blank sample oxidized at 800 ◦ C reflected after 210 cycles, which was due to the extensive spallation of the oxide scale. After spallation, fresh surfaces of the alloy were exposed to the atmosphere, hence, the subsequent oxidation rate increased. On the other hand, the oxide scale formed on the coated sample did not spall during the test period resulting in a smooth curve. The inflexion of the mass gain curve of the blank sample oxidized at 900 ◦ C occurred earlier than that at 800 ◦ C, the oxide scale on the blank sample spalled more easily. 3.3. Oxide phase Fig. 4 presents the phase composition from XRD analysis of the oxide scales formed on the specimens oxidized in air for 20 h at 900 and 1000 ◦ C. TiO2 (rutile) and ␣-Al2 O3 oxides could be detected in all the scales, but the relative intensity of the ␣-Al2 O3 peaks on the coated samples (Fig. 4(b) and (d)) was stronger than that on the uncoated samples (Fig. 4(a) and (c)). The examination reveals that more alumina formed on the coated samples during the oxidation processing. 3.4. Oxide morphologies Surface morphologies of the scales formed on samples oxidized isothermally for 20 h are shown in Fig. 5. After 20 h oxidation at 900 and 1000 ◦ C, the oxide scales formed on the blank samples (Fig. 5(a) and (c)) were consisted of equiaxed oxide grains by contrast to flake like oxide grains on the coated samples (Fig. 5(b) and (d)). After the oxidation at 1000 ◦ C, cracking and spallation of the oxide scale could be observed on the blank sample, but not on the coated sample. One problem with heterogeneous TiO2 /Al2 O3 scales is that they do
Fig. 3. Cyclic oxidation kinetics of blank and coated Ti3 Al based alloys at (a) 800 ◦ C and (b) 900 ◦ C in air.
180
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
Fig. 4. XRD patterns of the samples oxidized in air for 20 h. (a) Blank sample at 900 ◦ C, (b) coated sample at 900 ◦ C, (c) blank sample at 1000 ◦ C and (d) coated sample at 1000 ◦ C.
not adhere well. Cracking and spallation take place when the thickness of the oxide scale exceeded 10 m [22]. They spall during the oxidation process and especially during cooling process. The surface morphologies of the samples after the cyclic oxidation are not presented because they were similar to those after the isothermal oxidation. The cross-section morphologies of the samples oxidized isothermally for 20 h are shown in Fig. 6. It can be seen that the thickness of the scale formed on the coated sample (Fig. 6(b))
was thinner than that formed on the blank sample after the oxidation at 900 ◦ C (Fig. 6(a)). Because most of the oxide scale formed on the blank sample spalled during cooling after the oxidation at 1000 ◦ C, it was not possible to compare the total oxide thicknesses of the coated and uncoated samples based on the cross-section investigations. We noted that a white layer existed between the scale and the substrate on all the samples after oxidation (see in Fig. 6). The thickness of the white layer formed on the coated samples was much thinner than that of the blank samples, and the thickness of the white layer increased with increasing oxidation temperature. It seemed that the thickness of the white layers correlated with the degree of oxidation of the samples. In the area of the substrate alloy adjacent to the white layer, depletion of the white phase in the alloy was observed. With increasing oxidation temperature, the depletion became more obvious (Fig. 6(a) and (c)). Under the same oxidation condition, the white phase depleted less obviously in the coated alloy than that in the blank alloy (Fig. 6(c) and (d)). The concentration profiles along the depth in the crosssection of the coated sample oxidized isothermally at 1000 ◦ C for 20 h and circularly at 800 ◦ C for 300 cycles are shown in Figs. 7 and 8, respectively. It can be seen from Fig. 7 that the outer layer of the oxide scale was rich in Al and the inner layer was rich in Ti and the white layer was rich in Nb. As to the coated sample oxidized circularly at 800 ◦ C for 300 times (see in Fig. 8), the oxide scale mainly composed of Al2 O3 , and the white layer was nearly disappeared.
Fig. 5. Surface morphologies after isothermal oxidation for 20 h. Blank sample at 900 ◦ C (a) and 1000 ◦ C (c), coated sample at 900 ◦ C (b) and 1000 ◦ C (d).
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
181
Fig. 6. Cross-sections after isothermal oxidation for 20 h. Blank sample at 900 ◦ C (a) and 1000 ◦ C (c), coated sample at 900 ◦ C (b) and 1000 ◦ C (d).
4. Discussion The oxidation process of the Ti3 Al based alloys has been described in detail in Ref. [22]. It can be schematic illustrated in Fig. 9, both Al2 O3 and TiO2 formed at the initial stage, Al2 O3 formed more preferably than TiO2 . However, it did not form a continuous protective Al2 O3 layer because of the low activity of Al in the alloy (Fig. 9(a)). With the oxidation process continuing, oxygen diffused inwardly to react with titanium, and rutile formed beneath the non-continuous Al2 O3 layer. Part of the Al2 O3 was embedded in the rutile while others remained at the top of the scale (Fig. 9(b)). The Al2 O3 coating prepared in the present work was very thin (∼300 nm), so the effect of this Al2 O3 coating as a diffusion barrier was very limited. Under standard condition, the values of the Gibbs formation energy for Ti2 O3 (−1433.824 kJ/mol), which transforms to TiO2 by further oxidation, and Al2 O3 (−1582.271 kJ/mol) in the reaction of pure Ti and Al with oxygen respectively are quite close [23]. However, the equilibrium oxygen partial pressure of TiO2 is higher than that of Al2 O3 , so low oxygen pressure favors the formation of Al2 O3 . This conclusion is supported by the result that thermal treatment in an environment of low oxygen partial pressure improved the oxi-
dation resistance of Ti–Al based alloys [24]. The applied Al2 O3 coating could act to lower the oxygen partial pressure between the film and the substrate. It was beneficial to the formation of an Al2 O3 scale on the surface of the substrate, which was confirmed by the result shown in Fig. 8, where a continuous Al2 O3 layer formed on the coated alloy after long time cyclic oxidation at 800 ◦ C. Although this complete Al2 O3 scale did not form when the coated alloy was oxidized at 1000 ◦ C, its content in the mixed oxides still increased. As a result, the oxidation rate of the coated alloy was decreased, although not as dramatically as that at 800 ◦ C. EDS analysis in Fig. 7 shows that the white layers between the substrate and the oxide scale were rich in Nb and Al. Jiang et al. [5] studied the effect of Nb on the high-temperature oxidation of Ti- (0–50 at.%) Al, they also found an Nb and Al enriched layer in the outer layer of Ti–Al–Nb alloys adjacent to the scale/substrate interface. During the oxidation process, Ti, Al and Nb all diffused outward, but Ti and Al were oxidized preferentially because of their high oxygen affinity. Therefore, as the interface retracted continually into the inside of the alloy with the oxidation process, the alloy ahead of the interface became more and more enriched with Nb. The enrichment of Al at the oxide/alloy interface (seen in Fig. 7) was
182
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
Fig. 7. Concentration profiles along the depth in cross-section of the coated sample oxidized at 1000 ◦ C for 20 h.
thought to result from the depletion of Ti from its oxidation [5]. The more severe the oxidation, the more Ti consumed and the more distinct was the enrichment of Nb and Al (Fig. 6(a) and (c)). In the case of the Al2 O3 thin film coated samples (Fig. 6(b) and (d)), the oxidation rate of the Ti3 Al based alloy was reduced (Fig. 2), consequently the enrichment of Nb and Al was less. EDS analysis (Fig. 6) demonstrated that the white phase in the substrate was rich in Nb. Hence, the depletion of
Fig. 8. Concentration profiles along the depth in cross-section of the coated sample oxidized cyclically at 800 ◦ C for 300 times.
the white phase was related to the formation of the Nb enriched layer. The result of EDS analysis (Fig. 7) also showed that the content of niobium decreased gradually from the substrate/scale interface to the oxide/air interface. However, no niobium oxides were detected in the scales by XRD. Some of the niobium probably existed as solid solution in TiO2 . The defect structure of rutile consists of both titanium interstitials and oxygen vacancies, and the growth of titanium oxide in the scale proceeds by the inward
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
183
the test temperatures, especially during the cyclic oxidation at 800 ◦ C. This beneficial effect became weak after long time oxidation at 1000 ◦ C. 2. The applied thin film promoted the formation of Al2 O3 on the alloy increased the content of Al2 O3 in the mixed oxides, which caused the oxidation rate of the alloy to reduce. 3. An Nb and Al enrich layer was observed between the oxide scale and the substrate alloy for all of the samples, and the surface thin film could reduce the formation of this Nb and Al enrich layer through decreasing the oxidation rate of the alloys. Acknowledgements The authors are grateful for financial support from the National Science Foundation of China under grant No. 50371095. We also thank Professor R. Yang (IMR, CAS) for his affording the alloy. Fig. 9. The schematic illustration of the oxidation process of the Ti3 Al based alloys.
diffusion of oxygen and the outward diffusion of titanium [25], and the oxygen diffusion is the key step controlling the oxidation rate [26]. The substitution of Nb5+ for Ti4+ in rutile should decrease the concentration of vacancies, which would slow the diffusion of oxygen [5], thus suppressed the growth of rutile. The defect chemical equation can be expressed as, TiO2 → Ti2 O3 + 21 O2 ↑
(low PO2 )
(1)
References [1] [2] [3] [4] [5] [6] [7]
(2)
[8] [9] [10]
(3)
[11]
But the solubility of Nb in TiO2 is limited; Nb would exist in the scale as niobium oxide when more niobium is incorporated into the scale, causing the oxidation rate of the alloy to increase because of the high diffusion rate of oxygen in niobium oxide. Several researchers [5,27] have studied the effect of niobium content on the oxidation resistance of Ti–Al–Nb alloys. They suggested that the maximum oxidation resistance is obtained when the content of Nb is in the range of approximately 10–15 (at.%); both lower and higher Nb levels result in higher oxidation rates.
[12]
TiO2
Ti2 O3 −→2TiTi TiO2
+ V··O
+ 3O× O
1 Nb2 O5 −→2Nb·Ti + 2e + 4O× O + 2 O2 ↑
5. Conclusions Sol–gel derived Al2 O3 thin film has been applied on the surface of Ti–24Al–14Nb–3V–0.5Mo–0.3Si alloy. By investigating the isothermal and cyclic oxidation behaviors of the uncoated and coated alloys in air at 800–1000 ◦ C, the following conclusions can be drawn:
[13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
[23] [24] [25] [26]
1. The Al2 O3 thin film had a positive effect on the oxidation resistance of the alloy. The oxidation rates were reduced at
[27]
F.H. Froes, C. Suryanarayana, D. Eliezer, J. Mater. Sci. 27 (1992) 5113. C.M. Austin, C. Opin, Solid State Mater. Sci. 4 (1999) 239. E.A. Loria, Intermetallics 8 (2000) 1339. A. Rahmel, P.J. Spencer, Oxide Met. 35 (1991) 53. H.R. Jiang, M. Hirohasi, Y. Lu, H. Imanari, Scripta Mater. 46 (2002) 639. J.S. Wu, L.T. Zhang, F. Wang, K. Jiang, G.H. Qiu, Intermetallics 8 (2000) 19. R.A. Perkins, K.T. Chiang, G.H. Meier, Scripta Metall. Mater. 21 (1987) 1505. J. Subramahnyam, J. Mater. Sci. 23 (1988) 1906. Y. Shida, H. Anada, Oxid. Met. 45 (1996) 197. S. Becker, A. Rahemel, M. Schorr, M. Schutze, Oxid. Met. 38 (1992) 425. V. Gauthier, F. dettenwanger, M. Schutze, V. Shemet, W.J. Quadakkers, Oxid. Met. 59 (2003) 233. T. Nishimoto, T. Izumi, S. Hayashi, T. Narita, Intermetallics 11 (2003) 225. S.K. Jha, A.S. Khanna, C.S. Harendranath, Oxid. Met. 47 (1997) 465. L. Levin, A. Ginzburg, L. klinger, T. Werber, A. Katsman, P. Schaaf, Surf. Coat. Technol. 106 (1998) 209. S.-C. Kuang, Oxid. Met. 34 (1990) 217. Z.W. Li, W. Gao, M. Yoshihara, Y.D. He, Mater. Sci. Eng. 347A (2003) 243. C. Zeng, W. Wu, S. Zhu, J. Aero. Mater. 14 (1994) 21 (in Chinese). M.S. Chu, S.K. Wu, Surf. Coat. Technol. 179 (2004) 257. H. Li, K. Liang, L. Mei, S. Gu, S. Wang, Mater. Lett. 51 (2001) 320. S. Zhang, W.E. Lee, J. Eur. Ceram. Soc. 23 (2003) 1215. B.E. Yoldas, Ceram. Bull. 54 (1975) 289. G. Welsch, J.L. Smialek, J. Waldman, N.S. Jacobson, in: G. Welsch, P.D. Desai (Eds.), Oxidation and Corrsion of Intermetallic Alloys, West Lafayette, Indiana, 1996, p. 173. I. Barin, Thermochemical Data of Pure Substrates, VCH, Weinheim, 1989. E. Kobayashi, M. Yoshihara, R. Tanaka, High Temp. Technol. 8 (3) (1990) 179. P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, London and New York, 1988, p. 293. P. Kofstad, High Temperature Oxidation of Metals, Wiley, New York, 1996, p. 175. C. Leyens, H. Gedanitz, Scripta Metall. 41 (1999) 901.
Influence of sol–gel derived Al2O3 film on the oxidation behavior of a Ti3Al based alloy Ming Zhu a,b , Meishuan Li a,∗ , Yali Li a , Yanchun Zhou a a
High-Performance Ceramic Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of Chinese Academy of Sciences, Beijing 100039, China Received in revised form 6 September 2005; accepted 20 September 2005
Abstract Al2 O3 thin films were deposited on a Ti3 Al based alloy (Ti–24Al–14Nb–3V–0.5Mo–0.3Si) by sol–gel processing. Isothermal oxidation at temperatures of 900–1000 ◦ C and cyclic oxidation at 800–900 ◦ C were performed to test their effect on the oxidation behavior of the alloy. Results of the oxidation tests show that the oxidation parabolic rate constants of the alloy were reduced due to the applied thin film. This beneficial effect became weaker after longer oxidation time at 1000 ◦ C. TiO2 and Al2 O3 were the main phases formed on the alloy. The thin film could promote the growth of Al2 O3 , causing an increase of the Al2 O3 content in the composite oxides, sequentially decreased the oxidation rate. Nb/Al enriched as a layer in the alloy adjacent to the oxide/alloy interface in both the coated and uncoated alloy. The coated thin film decreased the thickness of the Nb/Al enrichment layer by reducing the scale growth rate. © 2005 Elsevier B.V. All rights reserved. Keywords: Ti3 Al based alloy; Sol–gel processing; Al2 O3 thin film; High-temperature oxidation
1. Introduction Ti3 Al based alloys are prospective high-temperature structural materials for aeroengine and automobile parts because of their excellent mechanical properties such as high specific strength and high elastic modulus at elevated temperatures [1–3]. However, poor oxidation resistance of these materials at high temperatures, which is attributed to the formation of mixed oxides containing the less protective TiO2 (rutile), and ␣-Al2 O3 , rather than a continuous protective ␣-Al2 O3 scale in oxidizing environments, limits their industrial applications. Rahmel and Spencer [4] have calculated the Al- and Ti-activities (aAl and aTi ) in Ti–Al intermetallic compounds. Their results show that aluminum concentration of over 50 at.% is needed for the growth of a single phase external Al2 O3 layer. Additions of some transition elements, such as Nb [5–8], Ta, Cr [9], Mo, W and Re can increase the activity ratio (aAl /aTi ) in these alloys, causing the protective alumina scale to be formed even when the aluminum concentration is less than 50 at.% [10]. However, some of the
∗
Corresponding author. Tel.: +86 24 23915913; fax: +86 24 23891320. E-mail address: [email protected] (M. Li).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.09.078
added elements may have negative effects on the mechanical properties of the alloys. This problem may be solved by depositing coatings on the alloys because the application of coatings may improve their oxidation resistance without degrading their mechanical properties. Many kinds of the coatings have been prepared in order to improve the oxidation resistance of Ti3 Al based alloys, such as diffusing coatings [11–15] and depositing coatings [16–18]. All of these coatings are meant to form a continuous protective Al2 O3 or Cr2 O3 layer on the surface during oxidation. Aluminide coatings are quite effective to improve the oxidation resistance of Ti3 Al based alloys, however, cracks generate easily in the coatings [15] due to large thermal stresses resulting from the thermal expansion mismatch between the coatings and the substrate alloys. Most of the deposited coatings usually have shortcomings of poor adhesion [17] and significant inter diffusion [18] between the coating and the substrate. Thus, new technologies or new coating systems need to be developed in order to further improve the oxidation resistance of Ti–Al based alloys. Sol–gel processing is widely used to prepare various functional thin films in which high-temperature oxidation resistant coatings are included [19,20], but its use as a protective coating on Ti–Al based alloys has not been reported. The generation of
178
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
cracks during thermal densification of the gel at around 500 ◦ C, and the lack of adhesion between these films and metals limit the use of this technique in preparing thick high-temperature oxidation resistant coatings, but cracking is seldom a problem for thin films. The aim of the present work is to study the effect of sol–gel derived Al2 O3 thin film on the selective oxidation of a Ti3 Al based alloy (Ti–24Al–14Nb–3V–0.5Mo–0.3Si, at.%), and to probe into the feasibility of using this sol–gel derived thin film as an oxidation resistant coating on this alloy.
reached the desired value. Cyclic oxidation was conducted at 800 ◦ C and 900 ◦ C. The specimens were kept in the furnace at 800 or 900 ◦ C for 1 h and then taken out to cool for 10 min in air. This process is defined as one cycle, and the cycle was repeated a few 100 times. An analytical balance with a sensitivity of 10−5 g was used to measure the weight of the specimens at intervals during the cyclic oxidation test. The surface morphologies and cross-sections of the oxidized samples were characterized by SEM equipped with EDS, while the phases of the oxide scales were analyzed using an X-ray diffractometer with Cu K␣ radiation.
2. Experimental 3. Results The nominal composition of the Ti3 Al based alloy is Ti–24Al–14Nb–3V–0.5Mo–0.3Si (at.%). The alloy ingot was prepared by melting twice in a consumable-electrode arc furnace under vacuum. Billets were cut from the ingots and, after descaling, hot press forged first in the -phase region at 1200 ◦ C and then in the (␣2 + 2 ) two-phase region at 1050 ◦ C with intermediate annealing. The heat treatment schedule of the alloy after forging was 1000 ◦ C/1 h/AC + 850 ◦ C/2 h/AC (AC stands for air-cooling). The alloy consists of three phases (␣2 + 2 + O). Specimens were cut into dimensions of 10 mm × 12 mm × 2 mm by using a spark wire machine, and their surfaces were grounded to 600 grit SiC paper. In order to improve the adhesion between the thin film and the substrate, the surfaces of the substrate were treated by sandblasting with steel abrasive powders. Then the samples were degreased ultrasonically in acetone, cleaned by distilled water and dried in air. Al2 O3 sol was prepared through the method used by Yoldas [21]. 20.4 g aluminum isopropoxide (Al(OC3 H7 )3 ), purity >98%, was hydrated in 270 ml distilled water at about 80 ◦ C with stirring for 1 h, then HNO3 was added into the solution to adjust the pH to about 3. The vessel containing the solution was kept constantly at 90 ◦ C for 16 h; a transparent sol was obtained. Samples were dipped into the sol and withdrawn at a rate between 3 and 30 cm/min, then dried in air for 15 min at room temperature. Thicker films could be obtained by repeating the dipping step. In this work samples were dipped and dried for five times. Thermal treatment was carried out in a tube furnace. The tube was first evacuated to a pressure of 2 × 10−1 Pa and then argon was introduced as a protective gas at a flow rate of 0.5 l/min. The tube was heated up to 500 ◦ C at the rate of 4 ◦ C/min and kept at this temperature for half an hour. Then it was heat up to 1000 ◦ C at the same heating rate and kept for 30 min. The target sol–gel Al2 O3 thin films were finally obtained by cooling the specimens to ambient temperature in the furnace. In order to understand the structure transformation of the gel during the annealing process, the dry gel was heated at 500, 700, 900 and 1000 ◦ C for half an hour, respectively, then grinded into powders and sieved. XRD powder analysis was used to examine the structure of the annealed gels. The continuous oxidation weight gain measurements were performed in a Sytses16/18(Setaram, France) microbalance at temperatures of 900–1000 ◦ C in air for 20 h. The samples were heated at a rate of 40 ◦ C/min to the test temperature. The mass change was recorded as a function of time once the temperature
3.1. Characteristics of the sol XRD patterns of the Al2 O3 gels treated for half an hour at 500–1000 ◦ C and then grinded into powers are presented in Fig. 1. Three main peaks of ␥-Al2 O3 were detected on the samples annealed at 500 and 700 ◦ C. The peaks corresponding to ␦-Al2 O3 also appeared after the treatment at 900 ◦ C. The gels treated below 1000 ◦ C was not well crystallized. The gel treated at 1000 ◦ C composed of ␣-Al2 O3 and a small mount of -Al2 O3 . Since a protective ␣-Al2 O3 layer is expected to provide the oxidation protection, the annealing temperature of the coating cannot be lower than 1000 ◦ C. But considering that the substrate alloy was heat treated at the temperature of 1000 ◦ C, the final annealing temperature of the coatings was better not to above 1000 ◦ C. 3.2. Oxidation kinetics The isothermal oxidation kinetics at different temperatures of the blank and coated samples are depicted in Fig. 2 as a function of the square of mass gain versus oxidation time. It can be seen that at the same oxidation temperature, the weight gains of the Al2 O3 coated samples were lower than those of Ti3 Al based alloy. The oxidation rate of the coated sample at 1000 ◦ C
Fig. 1. XRD patterns of the Al2 O3 gels after drying at (a) 500 ◦ C, (b) 700 ◦ C, (c) 900 ◦ C, and (d) 1000 ◦ C for 30 min.
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
Fig. 2. Plots of square of mass gain vs. time during the isothermal oxidation of blank and coated Ti3 Al based alloy at 900 and 1000 ◦ C. Table 1 Parabolic rate constants for different samples oxidized at 900 and 1000 ◦ C Parabolic rate constant, kp (mg2 cm−4 s−1 ) 900 ◦ C
1000 ◦ C
Blank alloy
2.20 × 10−4
8.32 × 10−4 (0–4 h) 3.96 × 10−4 (4–20 h)
Coated alloy
2.68 × 10−5
1.67 × 10−4
was even slower than that of the blank sample at 900 ◦ C. The oxidation kinetic of the blank sample at 1000 ◦ C deviated from the parabolic law while the other three curves presented in Fig. 2 followed the parabolic law after the first few hours. The parabolic rate constant (kp ) could be obtained by a linear fit of m2 = A + kp t, where A is a constant and the kp values are listed in Table 1. It is apparent that when coated with Al2 O3 film, the parabolic rate constants of the alloy decreased about 88% at 900 ◦ C, and roughly 60% at 1000 ◦ C. Fig. 3 shows the cyclic oxidation kinetics of the blank and coated samples at 800 ◦ C (a) and 900 ◦ C (b). The reason the cyclic oxidation was carried out at 800 and 900 ◦ C was because
179
the temperature at which Ti3 Al based alloys used in practical applications are usually between 800 and 900 ◦ C. It can be seen that the mass gain of each sample generally kept increasing during the cyclic oxidation. The oxidation rate of the coated sample was much lower than that of the blank sample when cyclic oxidized at 800 ◦ C (Fig. 3(a)). During the cyclic oxidation at 900 ◦ C (Fig. 3(b)), the sol–gel derived Al2 O3 thin film did not decrease the oxidation rate as remarkably as it did at 800 ◦ C. The curve corresponding to the blank sample oxidized at 800 ◦ C reflected after 210 cycles, which was due to the extensive spallation of the oxide scale. After spallation, fresh surfaces of the alloy were exposed to the atmosphere, hence, the subsequent oxidation rate increased. On the other hand, the oxide scale formed on the coated sample did not spall during the test period resulting in a smooth curve. The inflexion of the mass gain curve of the blank sample oxidized at 900 ◦ C occurred earlier than that at 800 ◦ C, the oxide scale on the blank sample spalled more easily. 3.3. Oxide phase Fig. 4 presents the phase composition from XRD analysis of the oxide scales formed on the specimens oxidized in air for 20 h at 900 and 1000 ◦ C. TiO2 (rutile) and ␣-Al2 O3 oxides could be detected in all the scales, but the relative intensity of the ␣-Al2 O3 peaks on the coated samples (Fig. 4(b) and (d)) was stronger than that on the uncoated samples (Fig. 4(a) and (c)). The examination reveals that more alumina formed on the coated samples during the oxidation processing. 3.4. Oxide morphologies Surface morphologies of the scales formed on samples oxidized isothermally for 20 h are shown in Fig. 5. After 20 h oxidation at 900 and 1000 ◦ C, the oxide scales formed on the blank samples (Fig. 5(a) and (c)) were consisted of equiaxed oxide grains by contrast to flake like oxide grains on the coated samples (Fig. 5(b) and (d)). After the oxidation at 1000 ◦ C, cracking and spallation of the oxide scale could be observed on the blank sample, but not on the coated sample. One problem with heterogeneous TiO2 /Al2 O3 scales is that they do
Fig. 3. Cyclic oxidation kinetics of blank and coated Ti3 Al based alloys at (a) 800 ◦ C and (b) 900 ◦ C in air.
180
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
Fig. 4. XRD patterns of the samples oxidized in air for 20 h. (a) Blank sample at 900 ◦ C, (b) coated sample at 900 ◦ C, (c) blank sample at 1000 ◦ C and (d) coated sample at 1000 ◦ C.
not adhere well. Cracking and spallation take place when the thickness of the oxide scale exceeded 10 m [22]. They spall during the oxidation process and especially during cooling process. The surface morphologies of the samples after the cyclic oxidation are not presented because they were similar to those after the isothermal oxidation. The cross-section morphologies of the samples oxidized isothermally for 20 h are shown in Fig. 6. It can be seen that the thickness of the scale formed on the coated sample (Fig. 6(b))
was thinner than that formed on the blank sample after the oxidation at 900 ◦ C (Fig. 6(a)). Because most of the oxide scale formed on the blank sample spalled during cooling after the oxidation at 1000 ◦ C, it was not possible to compare the total oxide thicknesses of the coated and uncoated samples based on the cross-section investigations. We noted that a white layer existed between the scale and the substrate on all the samples after oxidation (see in Fig. 6). The thickness of the white layer formed on the coated samples was much thinner than that of the blank samples, and the thickness of the white layer increased with increasing oxidation temperature. It seemed that the thickness of the white layers correlated with the degree of oxidation of the samples. In the area of the substrate alloy adjacent to the white layer, depletion of the white phase in the alloy was observed. With increasing oxidation temperature, the depletion became more obvious (Fig. 6(a) and (c)). Under the same oxidation condition, the white phase depleted less obviously in the coated alloy than that in the blank alloy (Fig. 6(c) and (d)). The concentration profiles along the depth in the crosssection of the coated sample oxidized isothermally at 1000 ◦ C for 20 h and circularly at 800 ◦ C for 300 cycles are shown in Figs. 7 and 8, respectively. It can be seen from Fig. 7 that the outer layer of the oxide scale was rich in Al and the inner layer was rich in Ti and the white layer was rich in Nb. As to the coated sample oxidized circularly at 800 ◦ C for 300 times (see in Fig. 8), the oxide scale mainly composed of Al2 O3 , and the white layer was nearly disappeared.
Fig. 5. Surface morphologies after isothermal oxidation for 20 h. Blank sample at 900 ◦ C (a) and 1000 ◦ C (c), coated sample at 900 ◦ C (b) and 1000 ◦ C (d).
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
181
Fig. 6. Cross-sections after isothermal oxidation for 20 h. Blank sample at 900 ◦ C (a) and 1000 ◦ C (c), coated sample at 900 ◦ C (b) and 1000 ◦ C (d).
4. Discussion The oxidation process of the Ti3 Al based alloys has been described in detail in Ref. [22]. It can be schematic illustrated in Fig. 9, both Al2 O3 and TiO2 formed at the initial stage, Al2 O3 formed more preferably than TiO2 . However, it did not form a continuous protective Al2 O3 layer because of the low activity of Al in the alloy (Fig. 9(a)). With the oxidation process continuing, oxygen diffused inwardly to react with titanium, and rutile formed beneath the non-continuous Al2 O3 layer. Part of the Al2 O3 was embedded in the rutile while others remained at the top of the scale (Fig. 9(b)). The Al2 O3 coating prepared in the present work was very thin (∼300 nm), so the effect of this Al2 O3 coating as a diffusion barrier was very limited. Under standard condition, the values of the Gibbs formation energy for Ti2 O3 (−1433.824 kJ/mol), which transforms to TiO2 by further oxidation, and Al2 O3 (−1582.271 kJ/mol) in the reaction of pure Ti and Al with oxygen respectively are quite close [23]. However, the equilibrium oxygen partial pressure of TiO2 is higher than that of Al2 O3 , so low oxygen pressure favors the formation of Al2 O3 . This conclusion is supported by the result that thermal treatment in an environment of low oxygen partial pressure improved the oxi-
dation resistance of Ti–Al based alloys [24]. The applied Al2 O3 coating could act to lower the oxygen partial pressure between the film and the substrate. It was beneficial to the formation of an Al2 O3 scale on the surface of the substrate, which was confirmed by the result shown in Fig. 8, where a continuous Al2 O3 layer formed on the coated alloy after long time cyclic oxidation at 800 ◦ C. Although this complete Al2 O3 scale did not form when the coated alloy was oxidized at 1000 ◦ C, its content in the mixed oxides still increased. As a result, the oxidation rate of the coated alloy was decreased, although not as dramatically as that at 800 ◦ C. EDS analysis in Fig. 7 shows that the white layers between the substrate and the oxide scale were rich in Nb and Al. Jiang et al. [5] studied the effect of Nb on the high-temperature oxidation of Ti- (0–50 at.%) Al, they also found an Nb and Al enriched layer in the outer layer of Ti–Al–Nb alloys adjacent to the scale/substrate interface. During the oxidation process, Ti, Al and Nb all diffused outward, but Ti and Al were oxidized preferentially because of their high oxygen affinity. Therefore, as the interface retracted continually into the inside of the alloy with the oxidation process, the alloy ahead of the interface became more and more enriched with Nb. The enrichment of Al at the oxide/alloy interface (seen in Fig. 7) was
182
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
Fig. 7. Concentration profiles along the depth in cross-section of the coated sample oxidized at 1000 ◦ C for 20 h.
thought to result from the depletion of Ti from its oxidation [5]. The more severe the oxidation, the more Ti consumed and the more distinct was the enrichment of Nb and Al (Fig. 6(a) and (c)). In the case of the Al2 O3 thin film coated samples (Fig. 6(b) and (d)), the oxidation rate of the Ti3 Al based alloy was reduced (Fig. 2), consequently the enrichment of Nb and Al was less. EDS analysis (Fig. 6) demonstrated that the white phase in the substrate was rich in Nb. Hence, the depletion of
Fig. 8. Concentration profiles along the depth in cross-section of the coated sample oxidized cyclically at 800 ◦ C for 300 times.
the white phase was related to the formation of the Nb enriched layer. The result of EDS analysis (Fig. 7) also showed that the content of niobium decreased gradually from the substrate/scale interface to the oxide/air interface. However, no niobium oxides were detected in the scales by XRD. Some of the niobium probably existed as solid solution in TiO2 . The defect structure of rutile consists of both titanium interstitials and oxygen vacancies, and the growth of titanium oxide in the scale proceeds by the inward
M. Zhu et al. / Materials Science and Engineering A 415 (2006) 177–183
183
the test temperatures, especially during the cyclic oxidation at 800 ◦ C. This beneficial effect became weak after long time oxidation at 1000 ◦ C. 2. The applied thin film promoted the formation of Al2 O3 on the alloy increased the content of Al2 O3 in the mixed oxides, which caused the oxidation rate of the alloy to reduce. 3. An Nb and Al enrich layer was observed between the oxide scale and the substrate alloy for all of the samples, and the surface thin film could reduce the formation of this Nb and Al enrich layer through decreasing the oxidation rate of the alloys. Acknowledgements The authors are grateful for financial support from the National Science Foundation of China under grant No. 50371095. We also thank Professor R. Yang (IMR, CAS) for his affording the alloy. Fig. 9. The schematic illustration of the oxidation process of the Ti3 Al based alloys.
diffusion of oxygen and the outward diffusion of titanium [25], and the oxygen diffusion is the key step controlling the oxidation rate [26]. The substitution of Nb5+ for Ti4+ in rutile should decrease the concentration of vacancies, which would slow the diffusion of oxygen [5], thus suppressed the growth of rutile. The defect chemical equation can be expressed as, TiO2 → Ti2 O3 + 21 O2 ↑
(low PO2 )
(1)
References [1] [2] [3] [4] [5] [6] [7]
(2)
[8] [9] [10]
(3)
[11]
But the solubility of Nb in TiO2 is limited; Nb would exist in the scale as niobium oxide when more niobium is incorporated into the scale, causing the oxidation rate of the alloy to increase because of the high diffusion rate of oxygen in niobium oxide. Several researchers [5,27] have studied the effect of niobium content on the oxidation resistance of Ti–Al–Nb alloys. They suggested that the maximum oxidation resistance is obtained when the content of Nb is in the range of approximately 10–15 (at.%); both lower and higher Nb levels result in higher oxidation rates.
[12]
TiO2
Ti2 O3 −→2TiTi TiO2
+ V··O
+ 3O× O
1 Nb2 O5 −→2Nb·Ti + 2e + 4O× O + 2 O2 ↑
5. Conclusions Sol–gel derived Al2 O3 thin film has been applied on the surface of Ti–24Al–14Nb–3V–0.5Mo–0.3Si alloy. By investigating the isothermal and cyclic oxidation behaviors of the uncoated and coated alloys in air at 800–1000 ◦ C, the following conclusions can be drawn:
[13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
[23] [24] [25] [26]
1. The Al2 O3 thin film had a positive effect on the oxidation resistance of the alloy. The oxidation rates were reduced at
[27]
F.H. Froes, C. Suryanarayana, D. Eliezer, J. Mater. Sci. 27 (1992) 5113. C.M. Austin, C. Opin, Solid State Mater. Sci. 4 (1999) 239. E.A. Loria, Intermetallics 8 (2000) 1339. A. Rahmel, P.J. Spencer, Oxide Met. 35 (1991) 53. H.R. Jiang, M. Hirohasi, Y. Lu, H. Imanari, Scripta Mater. 46 (2002) 639. J.S. Wu, L.T. Zhang, F. Wang, K. Jiang, G.H. Qiu, Intermetallics 8 (2000) 19. R.A. Perkins, K.T. Chiang, G.H. Meier, Scripta Metall. Mater. 21 (1987) 1505. J. Subramahnyam, J. Mater. Sci. 23 (1988) 1906. Y. Shida, H. Anada, Oxid. Met. 45 (1996) 197. S. Becker, A. Rahemel, M. Schorr, M. Schutze, Oxid. Met. 38 (1992) 425. V. Gauthier, F. dettenwanger, M. Schutze, V. Shemet, W.J. Quadakkers, Oxid. Met. 59 (2003) 233. T. Nishimoto, T. Izumi, S. Hayashi, T. Narita, Intermetallics 11 (2003) 225. S.K. Jha, A.S. Khanna, C.S. Harendranath, Oxid. Met. 47 (1997) 465. L. Levin, A. Ginzburg, L. klinger, T. Werber, A. Katsman, P. Schaaf, Surf. Coat. Technol. 106 (1998) 209. S.-C. Kuang, Oxid. Met. 34 (1990) 217. Z.W. Li, W. Gao, M. Yoshihara, Y.D. He, Mater. Sci. Eng. 347A (2003) 243. C. Zeng, W. Wu, S. Zhu, J. Aero. Mater. 14 (1994) 21 (in Chinese). M.S. Chu, S.K. Wu, Surf. Coat. Technol. 179 (2004) 257. H. Li, K. Liang, L. Mei, S. Gu, S. Wang, Mater. Lett. 51 (2001) 320. S. Zhang, W.E. Lee, J. Eur. Ceram. Soc. 23 (2003) 1215. B.E. Yoldas, Ceram. Bull. 54 (1975) 289. G. Welsch, J.L. Smialek, J. Waldman, N.S. Jacobson, in: G. Welsch, P.D. Desai (Eds.), Oxidation and Corrsion of Intermetallic Alloys, West Lafayette, Indiana, 1996, p. 173. I. Barin, Thermochemical Data of Pure Substrates, VCH, Weinheim, 1989. E. Kobayashi, M. Yoshihara, R. Tanaka, High Temp. Technol. 8 (3) (1990) 179. P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, London and New York, 1988, p. 293. P. Kofstad, High Temperature Oxidation of Metals, Wiley, New York, 1996, p. 175. C. Leyens, H. Gedanitz, Scripta Metall. 41 (1999) 901.