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sapphire interfaces: I. Structure and chemistry of interfaces
Materials Science and Engineering A349 (2003) 248 /255 www.elsevier.com/locate/msea
The effect of structure and chemistry on the strength of FeCrAl(Y)/ sapphire interfaces: I. Structure and chemistry of interfaces S.N. Basu a,*, H. Wu a,1, V. Gupta b, V. Kireev b a
Department of Manufacturing Engineering, Boston University, 15 St. Marys Street, Brookline, MA 02246, USA b Department of Mechanical and Aerospace Engineering, UCLA, Los Angeles, CA 90095, USA Received 21 May 2002; received in revised form 30 September 2002
Abstract FeCrAl/Al2O3 and FeCrAlY/Al2O3 interfaces were formed by sputter deposition of alloy thin films on A-plane sapphire substrates. The microstructure of both Y-containing and Y-free films consisted of single-phase BCC columnar grains of fairly uniform composition across the film thickness. Both films exhibited a high degree of 110 texture along the growth direction normal to the interface with no preferred in-plane orientation. In both cases, a thin amorphous layer was observed at the interface in the as-deposited state. Annealing the films for 16 h at 850 8C improved the crystalline quality of the film/substrate interface and led to grain growth within the films. The annealing did not lead to any significant changes in the film composition in the Y-free film. In the Y-containing film, yttrium oxide precipitates formed in the near interface region, depleting the adjacent alloy matrix of Y completely; accompanied by a significant enhancement of Al in the film, very close to the interface. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Reactive element effect; Structure of interfaces; Microstructure
1. Introduction The beneficial effect of reactive elements (‘reactive element’ or ‘rare earth’ effect) such as Y, La, Hf and Ce on oxidation behavior of alloys that form chromia or alumina scales has been well documented in the literature [1]. Extensive studies have been conducted to understand the significant improvement in adhesion of the oxide film to the reactive element-containing alloy substrate and many theories have been proposed to explain this phenomena. These include; modification of growth mechanism of the oxide scale, change in oxide morphology (suppression of oxide wrinkling), suppression of S segregation at the interface, improved chemical bonding at the interface, production of vacancy sinks, formation of oxide pegs at the interface, formation of a
* Corresponding author. Tel.: /1-617-353-6728; fax: /1-617-3535548. E-mail address: [email protected] (S.N. Basu). 1 Present address: Texas Instruments Inc., Attleboro, MA 02703, USA.
graded seal at the interface, and enhanced scale plasticity [1 /13]. Golightly et al. [2,3] studied oxide scale growth on Fe27Cr4Al alloys [14] with different amounts of Y doping (0, 0.023 and 0.82%, respectively). They reported that undoped-FeCrAl alloys developed laterally growing, equiaxed and highly convoluted oxide scales, resulting in extensive oxide detachment and subsequent spallation upon cooling. The lateral growth of the oxide was attributed to new oxide formation within the existing oxide layer due to simultaneous anion and cation diffusion, as proposed by Rhines and Wolf [4]. The growth of ‘wrinkles’ on thermally grown a-alumina scales on rare earth-free alloys has been recently studied by Tolpygo and Clarke [5,6], who showed that stress relaxation in the wrinkled scales was achieved by a mechanism of interfacial debonding, following by bucking and cracking of the detached region. In contrast, Ydoping led to oxide scales with columnar morphology in which lateral growth was suppressed, and the oxide/ alloy interface was flat and adherent. Tolpygo and Clarke showed a significant decrease in the lateral growth strain of a-alumina scales in the presence of Y
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 7 9 2 - X
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[5]. Study of oxidation mechanisms using two-stage O16/O-18 tracer oxidation techniques have resulted in conflicting results [7], with some studies indicating that the presence of Y suppressed Al diffusion [8], while others suggesting that it strongly enhanced Al diffusion [9]. However, in either case, transport of only one type of ion should eliminate lateral growth of the oxide. Another well studied theory of the reactive-element effect was proposed by Smeggil et al. [10], who proposed that Y getters the sulfur in the alloy and prevents the deleterious effects of S segregation at the oxide/alloy interface, thus leading to the enhanced scale adhesion. This hypothesis was supported by the observations of Smialek et al. [11], who reported that alumina scales on reactive element-free alloys that had been desulfurized by H2 annealing exhibited excellent scale adhesion. However, desulfurization by vacuum annealing does not lead to the same extent of improvement. This raised the question whether the improvement was the result of desulfurization itself or the effect of H2 and the reactiveelement on influencing interface strength by enhancing chemical bonding. Also, there is disagreement about whether S segregates to the metal/oxide interface [12] or only to the free metal surface [13]. In the former case, the segregation of S at the interface reduces the bonding strength; in the later case, it promotes the growth of existing voids at the interface by lowering surface energy. As seen by the above discussion, although significant research has been conducted to elucidate the mechanisms associated with the ‘reactive element effect’, there is no clear consensus on any one leading cause. It is entirely possible that the ‘reactive element effect’ occurs due to a combination of some of the above-discussed mechanisms. Interpretation of experimental data on thermally grown oxides is made more complicated by factors such as stresses in growing oxide scales, presence wavy interfaces with areas of stress concentrations, as well as the formation of microcracks. One factor that has not been directly studied is the effect of Y on the strength of the oxide/alloy interface. Jedlinki [7] has suggested that the increase of the interfacial bond strength allows the stress /relief mode in Y-containing alloys to be plastic deformation of the substrate, which is much less detrimental to scale adhesion than the mechanism of scale separation, cracking and spallation in Y-free alloys. This study is intended to carry out direct and quantitative measurements of the effect of Y on the interface strength of Al2O3/FeCrAl system. The emergence and maturation of laser spallation technique [15] make such measurements feasible. In this study, FeCrAl and FeCrAlY thin films were sputtered on sapphire substrates, instead of using the traditional method of generating the interface by alloy oxidation. Although the alloy/sapphire interfaces generated by this method may not be exactly
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identical to the alloy/alumina interfaces produced during high temperature oxidation of alloys, the complicating effects such as growth stress and the non-planar nature of interfaces formed during high temperature oxidation are avoided by using this method. Thus the comparison of the interfacial strengths of the Y-containing and Y-free samples allows us to concentrate solely on the effect of Y on the strength of the interface. Thus, the results of this study can be extended to interfaces produced during high temperature oxidation to address the issue of whether a change of intrinsic interfacial strength in the presence of Y is a contributing factor in the ‘rare earth effect’. This study combines electron microscopy and X-ray diffraction analysis with the laser spallation technique to investigate the effect of Y on the interface structure and strength. The structure of the FeCrAl(Y)/Al2O3 system is presented in this paper. The results of the study of the strength of FeCrAl(Y)/Al2O3 interfaces is presented in an accompanying paper [16].
2. Experimental procedure In this study, Y-containing and Y-free alloy thin films were sputtered on A-plane sapphire (single crystal Al2O3) substrates. The targets used for sputter deposition were vacuum cast by ACI Alloys Inc., using high purity elemental metals. The nominal compositions of the targets, listed in Table 1, were chosen such that they would be an alumina-former at elevated temperatures. The A-plane sapphire substrates, 10 /10 /0.5 mm in dimensions, were purchased with one side polished. Before deposition, the polished-side was chemically cleaned by sequentially washing and rinsing in trichloroethelene, acetone, methanol and DI water. Sputtering was carried out using the DC magnetron mode in a Denton Discovery 18 Deposition System. It has three magnetron sputter cathodes in confocal arrangement and rotating substrate table, supported by a 140 l s 1 turbomolecular pump that gives a presputter background vacuum of 10 7 Torr. The output power, the sputtering gas flow rate, the chamber pressure and the substrate temperature were systematically optimized for uniformity and adherence of the thin film. The optimal interface was achieved by using a Table 1 Average composition (wt.%) of the targets and as-deposited films Element
Fe Cr Al Y
FeCrAl
FeCrAlY
Target (nominal)
Film
Target (nominal)
Film
70 25 5 /
70.49 25.68 3.77 /
69 25 5 1
69.87 26.01 3.45 0.68
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combination of substrate heating to 285 8C, a chamber pressure of 6 /10 3 mbar and an output power of 220 W. The film thickness on the A-plane sapphire substrates after 1 h of sputtering was /0.85 mm. In order to study whether Y tended to segregate/ desegregate to the interface, some of the samples were annealed. For annealing, the samples were encapsulated in glass ampoules together with a Ta strip, which served as an oxygen getter. Prior to encapsulation, the ampoule was pumped down to 10 5 mbar and then back filled grade 5 Ar gas. The encapsulated samples were annealed at 850 8C for 16 h. The transmission electron microscopy (TEM) crosssectional samples were prepared following the standard sandwiching, sectioning, grinding, dimpling and ionmilling procedure. Ion-milling was carried out in a GATAN Duomill with a liquid N2 cold stage to minimize ion-milling artifacts. A JEM 2000FX electron microscope was used for TEM observations, and a JEM2010 electron microscope was used for HRTEM observations. Both microscopes were operated at 200KV. Analytical electron microscopy was carried out using a VG HB-603 dedicated STEM operated at 250 kV. The X-ray diffraction scans were conducted on a Rigaku 12 kW rotating anode X-ray diffractometer. The CuKa radiation is monochromated with a focusing graphite (002) monochromator crystal and collimated. The reflected X-ray is collected by rotating the scintillation detector to the appropriate angle. In this study, u / 2u scans, x-scans and f -scans were conducted. The schematic in Fig. 1 shows the appropriate angles [17]. For u /2u scan, the incoming X-ray beam is fixed, the sample and scintillation detector are rotated to u and 2u , respectively. This scan gives information on the d spacing of the diffracting planes in the samples. The film orientation quality was further examined by x-scans, in
Fig. 1. Schematic of X-ray diffraction geometry (partially adopted from [17]).
which the incoming X-ray beam and detector are fixed at u and 2u , respectively, while the sample is rotated to change x , as indicated in Fig. 1. Finally, the epitaxial relationship between film and substrate was examined by f-scans, in which the sample is rotated about its normal while u , 2u and x were held fixed at the peak position.
3. Results and discussion Examination of the as-deposited Y-containing and Yfree films in the SEM showed the film surface to be smooth and featureless. Examination of cross section in the TEM showed no significant differences in the film microstructures in the as-deposited Y-free and Y-containing films. Fig. 2 shows a cross-sectional bright field TEM micrograph of a FeCrAl film on A-plane sapphire substrate in the as-deposited state. The microstructure, consisting of columnar grains /30 nm in width was also typical of the FeCrAlY films. Composition analysis of cross-sectional samples in the STEM showed that the composition profiles were fairly uniform across the thickness of both Y-free and Ycontaining films, as evidenced by the composition profile across the FeCrAlY film plotted in Fig. 3. The average film compositions are listed in Table 1. As seen in Table 1, there was a slight enrichment of Cr and Fe and a more significant depletion of Al and Y in the films, as compared with the target. This is attributed to preferential sputtering of constituent elements of the target due to differences in equilibrium partial pressures and sticking coefficients. An important feature of the microstructure of both Yfree and Y-containing films in the as-deposited state was the presence of a very thin (/1 nm) amorphous interfacial layer, as shown in Fig. 4. This amorphous layer is probably a result of damage due to the bombardment of sputtered particles. Yuan et al. [18]
Fig. 2. Cross-sectional bright field TEM micrograph of an asdeposited FeCrAl film on A-plane sapphire substrate.
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Fig. 3. Composition profile across a FeCrAlY film in the as-deposited state.
have reported similar amorphous layers at Nb/sapphire interfaces formed by sputter deposition. The as-deposited and annealed films were examined by X-ray diffraction analysis. Fig. 5a shows a u /2u scan from an as-deposited Y-containing film. The figure shows a single {11 /20} diffraction peak at 37.98 from
Fig. 4. HREM micrograph of amorphous interfacial layer between the as-deposited FeCrAl film and A-plane sapphire substrate.
Fig. 5. u /2u X-ray diffraction scans of Y-containing film on A-plane sapphire in the a) as-deposited and b) annealed states, showing the {11 /20} sapphire peak at 37.98 and the {110} alloy peak at 44.98.
the A-plane sapphire substrate, consistent with a highquality single crystal. Interestingly, the scan also shows a single peak from the film at 44.98. This is consistent with a {110} peak from a BCC structure of the ferrous alloy. The presence of a single peak suggests a highly textured, single-phase BCC structure of the films. The texturing of the films was studied in detail by X-ray diffraction analysis, and a detailed description follows. A u /2u scan from an annealed Y-containing film, is presented in Fig. 5b, showing similar peaks in same locations. One obvious difference between the two spectra, taken under identical conditions in the X-ray diffractometer, is a fairly dramatic increase in the peak signal to the background noise ratio on annealing. One possible reason for this may be the crystallization of the interfacial region during annealing at 850 8C for 16 h. The presence of an amorphous interfacial region in the as-deposited films is due to damage by bombardment of sputtered ions. This loss of crystalline quality leads to an increase in the background noise and a decrease in the diffraction peak intensity. On annealing, this interfacial region crystallizes, and the improvement in the crystal quality is reflected in higher signal-to-noise ratio of the diffracted peaks. This hypothesis is confirmed by examining the interface of the annealed film by crosssectional TEM. Fig. 6 shows a high-resolution micrograph of the crystalline interface, showing that the substrate and film are crystalline right up to the interface. As discussed previously, the presence of a single diffraction peak from the film suggests film texture, with the 110 plane-normal oriented perpendicular to
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Fig. 7. x -scan of {200} FeCrAlY peak, showing a FWHM of the order of 28, consistent with a very high quality 110 texture.
Fig. 8. SAD pattern the interface of an annealed FeCrAl film showing that the (11 /20) sapphire diffraction spot lines up with the (110) film grain diffraction spot, consistent with a 110 growth texture of the film.
Fig. 6. HREM micrograph showing the crystalline interface between the FeCrAl film and the A-plane sapphire substrate after annealing at 850 8C for 16 h.
the interface. To study the quality of this 110 texture, a x-scan was carried out with the u value set at the Bragg angle for the {110} film peak. Fig. 7 shows such a x -scan for the annealed Y-containing film. The full width at half maximum (FWHM) of the {200} x -scan peak in this case is only of the order of 28, which corresponds to a very high quality 110 texture. The presence of excellent 110 texture was also observed in the TEM. Fig. 8 shows a SAD pattern from an interfacial region of a cross-sectional TEM sample of an annealed Y-free film. The SAD aperture was placed to select a single film grain. The figure shows that the (11 /20) diffraction spot of the sapphire substrate lines up with the (110) diffraction spot of the film, confirming
the 110 texture of the grains. Indexing of the diffraction spots indicates that the zone axis of the sapphire and the film grain are [0001] and [1 /13], respectively. Thus the orientation relation between the A-plane sapphire and this film grain can be expressed as: [0001]S½½[1 13]F (11 20)S½½(110)F;
(1)
where S and F stand for sapphire substrate and film, respectively. The grains oriented with their fastest growing 110 directions in the film growth direction soon dominate, leading to the growth of a 110 textured columnar microstructure. Having established that the columnar grains were oriented with their 110 axes being close to normal to the interface, the next obvious question is whether there is an in-plane orientation at the predominantly {11 /20} sapphire/{110} film interface. To study this, the {200} diffraction peak of the film was chosen. Fig. 9a shows a
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weak {200} diffraction peak at 65.38 in a u /2u scan from an annealed Y-containing film. To study the orientation of the 200 axes of the film grains, a x scan was carried out with the u value set at the Bragg angle for the {200} film diffraction peak (at ‘x’ in Fig. 9a). The scan, presented in Fig. 9b, shows no preferential orientation, indicating that the 200 axes of the film grains were randomly oriented. In order to confirm that the data in the x -scan did not correspond to the background noise, another x -scan was carried out with the u value set a few degrees off the Bragg angle for {200} (at ‘y ’ in Fig. 9a). This data, shown in Fig. 10b, does indeed measure the background noise. The average count intensity in Fig. 10a (/95 counts) is lower than that in Fig. 10b (/110 counts), indicating that the data did correspond to the weak (200) diffraction peak. The low intensity of this peak is consistent with the random distribution of the 200 film axis, thus spreading the diffraction intensity uniformly over 3608 of rotation. Annealing of the films was accompanied by grain growth to an average grain width of /90 nm for both Y-free and Y-containing films. In fact, the TEM and XRD studies on the as-deposited and annealed Y-free and Y-containing films showed very similar results in terms grain sizes in as-deposited an annealed states, as well as the crystallization of the amorphous interfacial region on annealing. However, one difference between
Fig. 9. (a) u /2u scan of annealed FeCrAlY film, showing a weak {200} diffraction peak at 65.38. (b) 8 -scans of the same sample, with the u value set at the {200} Bragg angle (x ) and a few degrees away from the {200} Bragg angle (y ).
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Fig. 10. Cross-sectional TEM bright-field micrograph of the annealed FeCrAlY film, showing the presence of a high-density of fine precipitates, some of which are marked by arrows.
the annealed Y-containing and Y-free films is depicted in Fig. 10, which shows a cross-sectional TEM micrograph of the annealed Y-containing film. The figure shows the presence of a high-density of fine precipitates in the annealed FeCrAlY film, some of which are marked by arrows. These precipitates were roughly in a 100 nm thick region of the film, adjacent to the interface. Fig. 11a shows a high-resolution micrograph of one of these precipitates, situated /10 nm from the interface. Since the EDX spectra from these precipitates showed the presence of yttrium and oxygen only, from here on these precipitates will be referred to as yttrium oxide precipitates. Fig. 11b shows an EDX spectrum from this precipitate, obtained using a 1 nm electron probe in the STEM. Also shown superimposed in Fig. 11b is an EDX spectrum from the matrix adjacent to the precipitate. A comparison of the spectra shows the precipitate to be an oxide that contains excess Y as compared with the matrix. The formation of yttrium oxide precipitates near the interface after annealing is believed to be due to diffusion of oxygen from the substrate to the film. The free energy change (per mole of oxygen) during oxidation for Y and Al are /240 and /210 kcal, respectively [19]. Thus, Y has a stronger tendency for oxidation than Al and is able to attract oxygen from the Al /O bond in sapphire to form yttrium oxide precipitates in the film close to the interface. The hypothesis is further supported by the STEM results in Table 2, which shows an Al concentration in excess of 35% in the film at a location of /5 nm from the interface. The oxygen thus
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a portion of an EDX spectrum from this precipitate superimposed with an EDX spectrum from the adjacent matrix, obtained using a 1 nm electron probe in the STEM. No Y signal can be seen in the EDX spectrum from the matrix in the figure. It appears that all the Y in the near-interface region of the annealed films has been tied up in the yttrium oxide precipitates. However, further away from the interface, there is no significant reduction in the Y content as compared with the asdeposited films. Secondly, the Al content in this very near-interface region is very high. It is conjectured that the source of this extra Al is the sapphire substrate, which dissociates to free up oxygen to form the more thermodynamically stable Y-rich oxide, allowing the freed Al to diffuse into the film. Also, the Al /O bonds at the sapphire substrate surface may be broken due to the bombardment of high energy sputtered ions during the initial stages of the deposition process, as well as from thermal energy at the high annealing temperature. Table 2 shows that there is an Al enrichment of the very near interface (/5 nm from interface) region of the Yfree FeCrAl film, although the extent of enrichment is much smaller as compared with the Y-containing film. Interestingly, no Y was found in the sapphire substrate for the FeCrAlY films, further supporting the fact that the near-interface Y was tied up as oxide and was not free to diffuse to the substrate.
Fig. 11. (a) High-resolution micrograph of a precipitate in an annealed FeCrAlY film situated /10 nm from the interface. (b) Y peak in the EDX spectrum from this precipitate superimposed on an EDX spectrum from the matrix adjacent to the precipitate, obtained using a 1 nm electron probe in the STEM.
released diffuses into the film and forms yttrium oxide precipitates in the film close to the interface. The composition of the matrix adjacent to the Y-rich oxide precipitate is listed in Table 2. This data reveals two very interesting phenomena. Firstly, the matrix of this FeCrAlY film, very close to the interface is virtually free of Y, or certainly the Y concentration is below the detectability limit of the VG603 STEM. Fig. 11b shows
Table 2 Near-interface (/5 nm away) film composition (wt.%) after annealing Element
FeCrAl
FeCrAlY (matrix adjacent to precipitate)
Fe Cr Al Y
69.68 25.18 5.18 /
48.32 18.5 36.62 /0
4. Conclusions In the present research, FeCrAl/Al2O3 and FeCrAlY/ Al2O3 interfaces were formed by sputtering deposition of alloy thin films on A-plane sapphire substrates. No significant differences in the film microstructure were observed between the Y-free and Y-containing films. In both cases, in the as-deposited state, a thin amorphous layer was observed at the interface, presumably due to damage during the sputter deposition process. The microstructure of both films consisted of columnar film grains of single phase BCC structure, with a high degree of 110 texture along the growth direction normal to the interface. The grains, however, had no preferred in-plane orientation relationships. The composition profiles of the constituent elements were found to be fairly uniform in the as-deposited films. Annealing the films for 16 h at 850 8C improved the crystalline quality of the film/substrate interface. The annealing did not lead to any significant changes in the composition of the Y-free film. In the Y-containing film, yttrium oxide precipitates formed in the near interface region, depleting the adjacent alloy matrix of Y completely; accompanied by a significant enhancement of Al in the film, very close to the interface.
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Acknowledgements This research was partially supported by the National Science Foundation (DMR 9705068). The authors would like to acknowledge the contributions of Professor Karl Ludwig, Dr Anthony Garratt-Reed and Michael Frongillo to this research. The electron microscopy studies were carried out at the Center for Electron Microscopy at MIT.
References [1] D.P. Whittle, J. Stringer, Phil. Trans. R. Soc. Lond. A295 (1980) 309. [2] F.A. Golightly, F.H. Stott, G.C. Wood, J. Electrochem. Soc. 126 (6) (1979) 1035. [3] F.A. Golightly, F.H. Stott, G.C. Wood, Oxid. Met. 14 (1980) 217. [4] F.N. Rhines, J.S. Wolf, Matall. Trans. 1 (1970) 1701. [5] V.K. Tolpygo, D.R. Clarke, Acta Mater. 46 (14) (1998) 5153.
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[6] V.K. Tolpygo, D.R. Clarke, Acta Mater. 46 (14) (1998) 5167. [7] J. Jedlinski, Solid State Phenomena 21/22 (1992) 335. [8] J. Jedlinski, S. Mrowec, Mater. Sci. Eng. 87 (1987) 281. [9] A. Glazkov, M. Gobel, J. Jedlinski, J. Schimmelpfennig, G. Borchardt, S. Weber, S. Scherrer, J.L. Coze, J. Phys. IV 5 (12) (1995) C7 /C381. [10] J.G. Smeggil, A.J. Shuskus, J. Vac. Sci. Technol. A4 (1986) 2577. [11] J.L. Smialek, B.K. Tubbs, Metall. Mater. Trans. 26A (2) (1995) 427. [12] P.Y. Hou, J. Stringer, Oxid. Met. 38 (11) (1992) 323. [13] H.J. Grabke, G. Kurbatov, H.J. Schmutzler, Oxid. Met. 43 (1995) 97. [14] All compositions are in weight% unless otherwise noted. [15] V. Gupta, J. Wu, A. Pronin, J. Am. Ceram. Soc. 80 (12) (1997) 3172. [16] S.N. Basu, H. Wu, V. Gupta, V. Kireev, ibid. [17] Figure is partially adopted from T. Lei, K.F. Ludwig Jr, T.D. Moustakas, J. Appl. Phys. 74 (1993) 4430. [18] J. Yuan, V. Gupta, M. Kim, Acta Metall. Mater. 43 (1995) 769. [19] T.B. Reed, Free Energy of Formation of Binary Compounds, MIT Press, Cambridge, MA, 1971.
The effect of structure and chemistry on the strength of FeCrAl(Y)/ sapphire interfaces: I. Structure and chemistry of interfaces S.N. Basu a,*, H. Wu a,1, V. Gupta b, V. Kireev b a
Department of Manufacturing Engineering, Boston University, 15 St. Marys Street, Brookline, MA 02246, USA b Department of Mechanical and Aerospace Engineering, UCLA, Los Angeles, CA 90095, USA Received 21 May 2002; received in revised form 30 September 2002
Abstract FeCrAl/Al2O3 and FeCrAlY/Al2O3 interfaces were formed by sputter deposition of alloy thin films on A-plane sapphire substrates. The microstructure of both Y-containing and Y-free films consisted of single-phase BCC columnar grains of fairly uniform composition across the film thickness. Both films exhibited a high degree of 110 texture along the growth direction normal to the interface with no preferred in-plane orientation. In both cases, a thin amorphous layer was observed at the interface in the as-deposited state. Annealing the films for 16 h at 850 8C improved the crystalline quality of the film/substrate interface and led to grain growth within the films. The annealing did not lead to any significant changes in the film composition in the Y-free film. In the Y-containing film, yttrium oxide precipitates formed in the near interface region, depleting the adjacent alloy matrix of Y completely; accompanied by a significant enhancement of Al in the film, very close to the interface. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Reactive element effect; Structure of interfaces; Microstructure
1. Introduction The beneficial effect of reactive elements (‘reactive element’ or ‘rare earth’ effect) such as Y, La, Hf and Ce on oxidation behavior of alloys that form chromia or alumina scales has been well documented in the literature [1]. Extensive studies have been conducted to understand the significant improvement in adhesion of the oxide film to the reactive element-containing alloy substrate and many theories have been proposed to explain this phenomena. These include; modification of growth mechanism of the oxide scale, change in oxide morphology (suppression of oxide wrinkling), suppression of S segregation at the interface, improved chemical bonding at the interface, production of vacancy sinks, formation of oxide pegs at the interface, formation of a
* Corresponding author. Tel.: /1-617-353-6728; fax: /1-617-3535548. E-mail address: [email protected] (S.N. Basu). 1 Present address: Texas Instruments Inc., Attleboro, MA 02703, USA.
graded seal at the interface, and enhanced scale plasticity [1 /13]. Golightly et al. [2,3] studied oxide scale growth on Fe27Cr4Al alloys [14] with different amounts of Y doping (0, 0.023 and 0.82%, respectively). They reported that undoped-FeCrAl alloys developed laterally growing, equiaxed and highly convoluted oxide scales, resulting in extensive oxide detachment and subsequent spallation upon cooling. The lateral growth of the oxide was attributed to new oxide formation within the existing oxide layer due to simultaneous anion and cation diffusion, as proposed by Rhines and Wolf [4]. The growth of ‘wrinkles’ on thermally grown a-alumina scales on rare earth-free alloys has been recently studied by Tolpygo and Clarke [5,6], who showed that stress relaxation in the wrinkled scales was achieved by a mechanism of interfacial debonding, following by bucking and cracking of the detached region. In contrast, Ydoping led to oxide scales with columnar morphology in which lateral growth was suppressed, and the oxide/ alloy interface was flat and adherent. Tolpygo and Clarke showed a significant decrease in the lateral growth strain of a-alumina scales in the presence of Y
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 7 9 2 - X
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[5]. Study of oxidation mechanisms using two-stage O16/O-18 tracer oxidation techniques have resulted in conflicting results [7], with some studies indicating that the presence of Y suppressed Al diffusion [8], while others suggesting that it strongly enhanced Al diffusion [9]. However, in either case, transport of only one type of ion should eliminate lateral growth of the oxide. Another well studied theory of the reactive-element effect was proposed by Smeggil et al. [10], who proposed that Y getters the sulfur in the alloy and prevents the deleterious effects of S segregation at the oxide/alloy interface, thus leading to the enhanced scale adhesion. This hypothesis was supported by the observations of Smialek et al. [11], who reported that alumina scales on reactive element-free alloys that had been desulfurized by H2 annealing exhibited excellent scale adhesion. However, desulfurization by vacuum annealing does not lead to the same extent of improvement. This raised the question whether the improvement was the result of desulfurization itself or the effect of H2 and the reactiveelement on influencing interface strength by enhancing chemical bonding. Also, there is disagreement about whether S segregates to the metal/oxide interface [12] or only to the free metal surface [13]. In the former case, the segregation of S at the interface reduces the bonding strength; in the later case, it promotes the growth of existing voids at the interface by lowering surface energy. As seen by the above discussion, although significant research has been conducted to elucidate the mechanisms associated with the ‘reactive element effect’, there is no clear consensus on any one leading cause. It is entirely possible that the ‘reactive element effect’ occurs due to a combination of some of the above-discussed mechanisms. Interpretation of experimental data on thermally grown oxides is made more complicated by factors such as stresses in growing oxide scales, presence wavy interfaces with areas of stress concentrations, as well as the formation of microcracks. One factor that has not been directly studied is the effect of Y on the strength of the oxide/alloy interface. Jedlinki [7] has suggested that the increase of the interfacial bond strength allows the stress /relief mode in Y-containing alloys to be plastic deformation of the substrate, which is much less detrimental to scale adhesion than the mechanism of scale separation, cracking and spallation in Y-free alloys. This study is intended to carry out direct and quantitative measurements of the effect of Y on the interface strength of Al2O3/FeCrAl system. The emergence and maturation of laser spallation technique [15] make such measurements feasible. In this study, FeCrAl and FeCrAlY thin films were sputtered on sapphire substrates, instead of using the traditional method of generating the interface by alloy oxidation. Although the alloy/sapphire interfaces generated by this method may not be exactly
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identical to the alloy/alumina interfaces produced during high temperature oxidation of alloys, the complicating effects such as growth stress and the non-planar nature of interfaces formed during high temperature oxidation are avoided by using this method. Thus the comparison of the interfacial strengths of the Y-containing and Y-free samples allows us to concentrate solely on the effect of Y on the strength of the interface. Thus, the results of this study can be extended to interfaces produced during high temperature oxidation to address the issue of whether a change of intrinsic interfacial strength in the presence of Y is a contributing factor in the ‘rare earth effect’. This study combines electron microscopy and X-ray diffraction analysis with the laser spallation technique to investigate the effect of Y on the interface structure and strength. The structure of the FeCrAl(Y)/Al2O3 system is presented in this paper. The results of the study of the strength of FeCrAl(Y)/Al2O3 interfaces is presented in an accompanying paper [16].
2. Experimental procedure In this study, Y-containing and Y-free alloy thin films were sputtered on A-plane sapphire (single crystal Al2O3) substrates. The targets used for sputter deposition were vacuum cast by ACI Alloys Inc., using high purity elemental metals. The nominal compositions of the targets, listed in Table 1, were chosen such that they would be an alumina-former at elevated temperatures. The A-plane sapphire substrates, 10 /10 /0.5 mm in dimensions, were purchased with one side polished. Before deposition, the polished-side was chemically cleaned by sequentially washing and rinsing in trichloroethelene, acetone, methanol and DI water. Sputtering was carried out using the DC magnetron mode in a Denton Discovery 18 Deposition System. It has three magnetron sputter cathodes in confocal arrangement and rotating substrate table, supported by a 140 l s 1 turbomolecular pump that gives a presputter background vacuum of 10 7 Torr. The output power, the sputtering gas flow rate, the chamber pressure and the substrate temperature were systematically optimized for uniformity and adherence of the thin film. The optimal interface was achieved by using a Table 1 Average composition (wt.%) of the targets and as-deposited films Element
Fe Cr Al Y
FeCrAl
FeCrAlY
Target (nominal)
Film
Target (nominal)
Film
70 25 5 /
70.49 25.68 3.77 /
69 25 5 1
69.87 26.01 3.45 0.68
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combination of substrate heating to 285 8C, a chamber pressure of 6 /10 3 mbar and an output power of 220 W. The film thickness on the A-plane sapphire substrates after 1 h of sputtering was /0.85 mm. In order to study whether Y tended to segregate/ desegregate to the interface, some of the samples were annealed. For annealing, the samples were encapsulated in glass ampoules together with a Ta strip, which served as an oxygen getter. Prior to encapsulation, the ampoule was pumped down to 10 5 mbar and then back filled grade 5 Ar gas. The encapsulated samples were annealed at 850 8C for 16 h. The transmission electron microscopy (TEM) crosssectional samples were prepared following the standard sandwiching, sectioning, grinding, dimpling and ionmilling procedure. Ion-milling was carried out in a GATAN Duomill with a liquid N2 cold stage to minimize ion-milling artifacts. A JEM 2000FX electron microscope was used for TEM observations, and a JEM2010 electron microscope was used for HRTEM observations. Both microscopes were operated at 200KV. Analytical electron microscopy was carried out using a VG HB-603 dedicated STEM operated at 250 kV. The X-ray diffraction scans were conducted on a Rigaku 12 kW rotating anode X-ray diffractometer. The CuKa radiation is monochromated with a focusing graphite (002) monochromator crystal and collimated. The reflected X-ray is collected by rotating the scintillation detector to the appropriate angle. In this study, u / 2u scans, x-scans and f -scans were conducted. The schematic in Fig. 1 shows the appropriate angles [17]. For u /2u scan, the incoming X-ray beam is fixed, the sample and scintillation detector are rotated to u and 2u , respectively. This scan gives information on the d spacing of the diffracting planes in the samples. The film orientation quality was further examined by x-scans, in
Fig. 1. Schematic of X-ray diffraction geometry (partially adopted from [17]).
which the incoming X-ray beam and detector are fixed at u and 2u , respectively, while the sample is rotated to change x , as indicated in Fig. 1. Finally, the epitaxial relationship between film and substrate was examined by f-scans, in which the sample is rotated about its normal while u , 2u and x were held fixed at the peak position.
3. Results and discussion Examination of the as-deposited Y-containing and Yfree films in the SEM showed the film surface to be smooth and featureless. Examination of cross section in the TEM showed no significant differences in the film microstructures in the as-deposited Y-free and Y-containing films. Fig. 2 shows a cross-sectional bright field TEM micrograph of a FeCrAl film on A-plane sapphire substrate in the as-deposited state. The microstructure, consisting of columnar grains /30 nm in width was also typical of the FeCrAlY films. Composition analysis of cross-sectional samples in the STEM showed that the composition profiles were fairly uniform across the thickness of both Y-free and Ycontaining films, as evidenced by the composition profile across the FeCrAlY film plotted in Fig. 3. The average film compositions are listed in Table 1. As seen in Table 1, there was a slight enrichment of Cr and Fe and a more significant depletion of Al and Y in the films, as compared with the target. This is attributed to preferential sputtering of constituent elements of the target due to differences in equilibrium partial pressures and sticking coefficients. An important feature of the microstructure of both Yfree and Y-containing films in the as-deposited state was the presence of a very thin (/1 nm) amorphous interfacial layer, as shown in Fig. 4. This amorphous layer is probably a result of damage due to the bombardment of sputtered particles. Yuan et al. [18]
Fig. 2. Cross-sectional bright field TEM micrograph of an asdeposited FeCrAl film on A-plane sapphire substrate.
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Fig. 3. Composition profile across a FeCrAlY film in the as-deposited state.
have reported similar amorphous layers at Nb/sapphire interfaces formed by sputter deposition. The as-deposited and annealed films were examined by X-ray diffraction analysis. Fig. 5a shows a u /2u scan from an as-deposited Y-containing film. The figure shows a single {11 /20} diffraction peak at 37.98 from
Fig. 4. HREM micrograph of amorphous interfacial layer between the as-deposited FeCrAl film and A-plane sapphire substrate.
Fig. 5. u /2u X-ray diffraction scans of Y-containing film on A-plane sapphire in the a) as-deposited and b) annealed states, showing the {11 /20} sapphire peak at 37.98 and the {110} alloy peak at 44.98.
the A-plane sapphire substrate, consistent with a highquality single crystal. Interestingly, the scan also shows a single peak from the film at 44.98. This is consistent with a {110} peak from a BCC structure of the ferrous alloy. The presence of a single peak suggests a highly textured, single-phase BCC structure of the films. The texturing of the films was studied in detail by X-ray diffraction analysis, and a detailed description follows. A u /2u scan from an annealed Y-containing film, is presented in Fig. 5b, showing similar peaks in same locations. One obvious difference between the two spectra, taken under identical conditions in the X-ray diffractometer, is a fairly dramatic increase in the peak signal to the background noise ratio on annealing. One possible reason for this may be the crystallization of the interfacial region during annealing at 850 8C for 16 h. The presence of an amorphous interfacial region in the as-deposited films is due to damage by bombardment of sputtered ions. This loss of crystalline quality leads to an increase in the background noise and a decrease in the diffraction peak intensity. On annealing, this interfacial region crystallizes, and the improvement in the crystal quality is reflected in higher signal-to-noise ratio of the diffracted peaks. This hypothesis is confirmed by examining the interface of the annealed film by crosssectional TEM. Fig. 6 shows a high-resolution micrograph of the crystalline interface, showing that the substrate and film are crystalline right up to the interface. As discussed previously, the presence of a single diffraction peak from the film suggests film texture, with the 110 plane-normal oriented perpendicular to
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Fig. 7. x -scan of {200} FeCrAlY peak, showing a FWHM of the order of 28, consistent with a very high quality 110 texture.
Fig. 8. SAD pattern the interface of an annealed FeCrAl film showing that the (11 /20) sapphire diffraction spot lines up with the (110) film grain diffraction spot, consistent with a 110 growth texture of the film.
Fig. 6. HREM micrograph showing the crystalline interface between the FeCrAl film and the A-plane sapphire substrate after annealing at 850 8C for 16 h.
the interface. To study the quality of this 110 texture, a x-scan was carried out with the u value set at the Bragg angle for the {110} film peak. Fig. 7 shows such a x -scan for the annealed Y-containing film. The full width at half maximum (FWHM) of the {200} x -scan peak in this case is only of the order of 28, which corresponds to a very high quality 110 texture. The presence of excellent 110 texture was also observed in the TEM. Fig. 8 shows a SAD pattern from an interfacial region of a cross-sectional TEM sample of an annealed Y-free film. The SAD aperture was placed to select a single film grain. The figure shows that the (11 /20) diffraction spot of the sapphire substrate lines up with the (110) diffraction spot of the film, confirming
the 110 texture of the grains. Indexing of the diffraction spots indicates that the zone axis of the sapphire and the film grain are [0001] and [1 /13], respectively. Thus the orientation relation between the A-plane sapphire and this film grain can be expressed as: [0001]S½½[1 13]F (11 20)S½½(110)F;
(1)
where S and F stand for sapphire substrate and film, respectively. The grains oriented with their fastest growing 110 directions in the film growth direction soon dominate, leading to the growth of a 110 textured columnar microstructure. Having established that the columnar grains were oriented with their 110 axes being close to normal to the interface, the next obvious question is whether there is an in-plane orientation at the predominantly {11 /20} sapphire/{110} film interface. To study this, the {200} diffraction peak of the film was chosen. Fig. 9a shows a
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weak {200} diffraction peak at 65.38 in a u /2u scan from an annealed Y-containing film. To study the orientation of the 200 axes of the film grains, a x scan was carried out with the u value set at the Bragg angle for the {200} film diffraction peak (at ‘x’ in Fig. 9a). The scan, presented in Fig. 9b, shows no preferential orientation, indicating that the 200 axes of the film grains were randomly oriented. In order to confirm that the data in the x -scan did not correspond to the background noise, another x -scan was carried out with the u value set a few degrees off the Bragg angle for {200} (at ‘y ’ in Fig. 9a). This data, shown in Fig. 10b, does indeed measure the background noise. The average count intensity in Fig. 10a (/95 counts) is lower than that in Fig. 10b (/110 counts), indicating that the data did correspond to the weak (200) diffraction peak. The low intensity of this peak is consistent with the random distribution of the 200 film axis, thus spreading the diffraction intensity uniformly over 3608 of rotation. Annealing of the films was accompanied by grain growth to an average grain width of /90 nm for both Y-free and Y-containing films. In fact, the TEM and XRD studies on the as-deposited and annealed Y-free and Y-containing films showed very similar results in terms grain sizes in as-deposited an annealed states, as well as the crystallization of the amorphous interfacial region on annealing. However, one difference between
Fig. 9. (a) u /2u scan of annealed FeCrAlY film, showing a weak {200} diffraction peak at 65.38. (b) 8 -scans of the same sample, with the u value set at the {200} Bragg angle (x ) and a few degrees away from the {200} Bragg angle (y ).
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Fig. 10. Cross-sectional TEM bright-field micrograph of the annealed FeCrAlY film, showing the presence of a high-density of fine precipitates, some of which are marked by arrows.
the annealed Y-containing and Y-free films is depicted in Fig. 10, which shows a cross-sectional TEM micrograph of the annealed Y-containing film. The figure shows the presence of a high-density of fine precipitates in the annealed FeCrAlY film, some of which are marked by arrows. These precipitates were roughly in a 100 nm thick region of the film, adjacent to the interface. Fig. 11a shows a high-resolution micrograph of one of these precipitates, situated /10 nm from the interface. Since the EDX spectra from these precipitates showed the presence of yttrium and oxygen only, from here on these precipitates will be referred to as yttrium oxide precipitates. Fig. 11b shows an EDX spectrum from this precipitate, obtained using a 1 nm electron probe in the STEM. Also shown superimposed in Fig. 11b is an EDX spectrum from the matrix adjacent to the precipitate. A comparison of the spectra shows the precipitate to be an oxide that contains excess Y as compared with the matrix. The formation of yttrium oxide precipitates near the interface after annealing is believed to be due to diffusion of oxygen from the substrate to the film. The free energy change (per mole of oxygen) during oxidation for Y and Al are /240 and /210 kcal, respectively [19]. Thus, Y has a stronger tendency for oxidation than Al and is able to attract oxygen from the Al /O bond in sapphire to form yttrium oxide precipitates in the film close to the interface. The hypothesis is further supported by the STEM results in Table 2, which shows an Al concentration in excess of 35% in the film at a location of /5 nm from the interface. The oxygen thus
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a portion of an EDX spectrum from this precipitate superimposed with an EDX spectrum from the adjacent matrix, obtained using a 1 nm electron probe in the STEM. No Y signal can be seen in the EDX spectrum from the matrix in the figure. It appears that all the Y in the near-interface region of the annealed films has been tied up in the yttrium oxide precipitates. However, further away from the interface, there is no significant reduction in the Y content as compared with the asdeposited films. Secondly, the Al content in this very near-interface region is very high. It is conjectured that the source of this extra Al is the sapphire substrate, which dissociates to free up oxygen to form the more thermodynamically stable Y-rich oxide, allowing the freed Al to diffuse into the film. Also, the Al /O bonds at the sapphire substrate surface may be broken due to the bombardment of high energy sputtered ions during the initial stages of the deposition process, as well as from thermal energy at the high annealing temperature. Table 2 shows that there is an Al enrichment of the very near interface (/5 nm from interface) region of the Yfree FeCrAl film, although the extent of enrichment is much smaller as compared with the Y-containing film. Interestingly, no Y was found in the sapphire substrate for the FeCrAlY films, further supporting the fact that the near-interface Y was tied up as oxide and was not free to diffuse to the substrate.
Fig. 11. (a) High-resolution micrograph of a precipitate in an annealed FeCrAlY film situated /10 nm from the interface. (b) Y peak in the EDX spectrum from this precipitate superimposed on an EDX spectrum from the matrix adjacent to the precipitate, obtained using a 1 nm electron probe in the STEM.
released diffuses into the film and forms yttrium oxide precipitates in the film close to the interface. The composition of the matrix adjacent to the Y-rich oxide precipitate is listed in Table 2. This data reveals two very interesting phenomena. Firstly, the matrix of this FeCrAlY film, very close to the interface is virtually free of Y, or certainly the Y concentration is below the detectability limit of the VG603 STEM. Fig. 11b shows
Table 2 Near-interface (/5 nm away) film composition (wt.%) after annealing Element
FeCrAl
FeCrAlY (matrix adjacent to precipitate)
Fe Cr Al Y
69.68 25.18 5.18 /
48.32 18.5 36.62 /0
4. Conclusions In the present research, FeCrAl/Al2O3 and FeCrAlY/ Al2O3 interfaces were formed by sputtering deposition of alloy thin films on A-plane sapphire substrates. No significant differences in the film microstructure were observed between the Y-free and Y-containing films. In both cases, in the as-deposited state, a thin amorphous layer was observed at the interface, presumably due to damage during the sputter deposition process. The microstructure of both films consisted of columnar film grains of single phase BCC structure, with a high degree of 110 texture along the growth direction normal to the interface. The grains, however, had no preferred in-plane orientation relationships. The composition profiles of the constituent elements were found to be fairly uniform in the as-deposited films. Annealing the films for 16 h at 850 8C improved the crystalline quality of the film/substrate interface. The annealing did not lead to any significant changes in the composition of the Y-free film. In the Y-containing film, yttrium oxide precipitates formed in the near interface region, depleting the adjacent alloy matrix of Y completely; accompanied by a significant enhancement of Al in the film, very close to the interface.
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Acknowledgements This research was partially supported by the National Science Foundation (DMR 9705068). The authors would like to acknowledge the contributions of Professor Karl Ludwig, Dr Anthony Garratt-Reed and Michael Frongillo to this research. The electron microscopy studies were carried out at the Center for Electron Microscopy at MIT.
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[6] V.K. Tolpygo, D.R. Clarke, Acta Mater. 46 (14) (1998) 5167. [7] J. Jedlinski, Solid State Phenomena 21/22 (1992) 335. [8] J. Jedlinski, S. Mrowec, Mater. Sci. Eng. 87 (1987) 281. [9] A. Glazkov, M. Gobel, J. Jedlinski, J. Schimmelpfennig, G. Borchardt, S. Weber, S. Scherrer, J.L. Coze, J. Phys. IV 5 (12) (1995) C7 /C381. [10] J.G. Smeggil, A.J. Shuskus, J. Vac. Sci. Technol. A4 (1986) 2577. [11] J.L. Smialek, B.K. Tubbs, Metall. Mater. Trans. 26A (2) (1995) 427. [12] P.Y. Hou, J. Stringer, Oxid. Met. 38 (11) (1992) 323. [13] H.J. Grabke, G. Kurbatov, H.J. Schmutzler, Oxid. Met. 43 (1995) 97. [14] All compositions are in weight% unless otherwise noted. [15] V. Gupta, J. Wu, A. Pronin, J. Am. Ceram. Soc. 80 (12) (1997) 3172. [16] S.N. Basu, H. Wu, V. Gupta, V. Kireev, ibid. [17] Figure is partially adopted from T. Lei, K.F. Ludwig Jr, T.D. Moustakas, J. Appl. Phys. 74 (1993) 4430. [18] J. Yuan, V. Gupta, M. Kim, Acta Metall. Mater. 43 (1995) 769. [19] T.B. Reed, Free Energy of Formation of Binary Compounds, MIT Press, Cambridge, MA, 1971.