- Email: [email protected]
Passivity and its breakdown in Al-based amorphous alloys
Materials Chemistry and Physics 92 (2005) 348–353
Passivity and its breakdown in Al-based amorphous alloys M. Janik-Czachor a,∗ , A. Jaskiewicz a,b , M. Dolata a,c , Z. Werner a,d a
Physical Chemistry of Materials Center, Institute of Physical Chemistry, Polish Academy of Sciences, 01-224 Warsaw, Kasprzaka str. 44/52, Poland b Faculty of Materials Science, Warsaw University of Technology, Warsaw, Poland c Warsaw Agriculture University, Warsaw, Poland d The Andrzej Soltan Institute for Nuclear Studies, Swierk, Poland Received 8 September 2004; received in revised form 29 October 2004; accepted 15 November 2004
Abstract The effect of Cl− ions (0–1 M NaCl) on the anodic behaviour of two series of sputter-deposited Al–Ta (11–46 at.%) and Al–Nb (13–46 at.%) amorphous alloys was investigated in a neutral electrolyte by using electrochemical methods and the results were compared with those for Al–Mo and Al–W alloys. In solutions containing Cl− the alloys broke down at and above the pit nucleation potential Enp . They exhibited a stable passivity within a much wider (E > 1 V) potential range than crystalline Al. The anodically formed oxide films in Cl− -free electrolyte were characterized by surface analytical measurements including high-energy resolution AES. Galvanostatic polarization was used to produce thicker anodic films than those formed under natural conditions, to facilitate detailed examination of their depth composition profiles. It was found that the anodic films grown on Al–Ta and Al–Nb alloys had a homogeneous structure in contrast to those on Al–Mo and Al–W. Transition metal Ta or Nb was detected at any depth of the composition profile. A relative enrichment of refractory metal within the inner part of the anodic film was not found, for these alloys. Tentative mechanisms of enhanced stability of the passive state of the Al-based amorphous alloys (AAs) are discussed. © 2005 Elsevier B.V. All rights reserved. Keywords: Amorphous alloys; Al-based alloys; Passivity breakdown; Pit nucleation potential; Insulating films; High-energy resolution AES analysis
1. Introduction Amorphous alloys (AAs) offer a new possibility of studying the effects of alloying elements on corrosion and passivity in a wide concentration range. Their great compositional flexibility, makes it possible to study the chemical effects without any interaction of structural effect [1–4]. One can, for example, produce homogeneous Al-based AAs containing from ∼11 at.% up to 50 at.% of a refractory metal (Mo, W, Ta, or Nb) [4], while in the crystalline state the solubility of these refractory metals is below 1 at.% [5]. Refractory metals are known for their beneficial effect on the passivity of many alloys, both crystalline and amorphous, although the mechanism of their action is still a subject of discussion [6–19]. AAs thus offer the opportunity of gaining an insight into the role of these elements, when studying ∗
Corresponding author. Tel.: +48 22 3433225; fax: +48 22 6325276. E-mail address: [email protected] (M. Janik-Czachor).
0254-0584/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2004.11.017
their effect also at higher concentrations. In particular, one can check whether the preferential dissolution of Al and/or different transport rates within the passive film of the components determines the anodic behavior [24], or whether the effect of refractory metal is confined to the action of “metal dissolution moderators” rather than to “passivity promoters”, as suggested by Marcus [7]. As the previous results [14,18] strongly suggest the former effect for Al–W and Al–Mo, it was thus crucial to study Al–Ta and Al–Nb AAs in order to prove the generality of the “metal dissolution moderators” model. While the elements considered, Mo, W, Ta, and Nb, belong to a group of refractory metals with a high energy of Me–Me bond [7] (the Me–Me bond energy for all the mentioned refractory metals is between 170 and 220 kJ mol−1 , while the corresponding value for Al is <60 kJ mol−1 [7]), their affinity to oxygen differs considerably from one another. Table 1 lists the corresponding values of free enthalpy of formation of an oxide (G◦ ). The G◦ value for Al2 O3 is also given.
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353 Table 1 Gibbs free enthalpy of formation of relevant oxide Oxide
G◦ (kJ mol−1 )
MoO3 WO3 Al2 O3 Nb2 O5 Ta2 O5
−667.5 −764.2 −1582.72 −1765.86 −1910.99
While G◦ for molybdenum and tungsten oxide is lower than that for Al2 O3 , the corresponding values for Ta2 O5 and Nb2 O5 are much higher. Thus, the anodic behavior of the Ta and Nb AAs may be different than that of the other refractory metals mentioned above. To carefully investigate the depth distribution of both components within the protective surface film, thick anodic films formed on Al–Mo and Al–W AAs were studied [13,14,18] as model systems for typical thin passivating films. The same procedure was adopted in this work to study anodic films on Al–Ta and Al–Nb AAs, to provide a proper basis for comparison. The aim of this work was to investigate the stability of the passive state of Al–Ta and Al–Nb AAs, and to compare the results with those for Al–Mo/W AAs in order to gain insight into the beneficial effect of the refractory metals.
2. Experimental The materials used in the investigations were Al–Ta and Al–Nb amorphous alloys obtained by sputter deposition onto Pyrex glass substrates. The exposed surface area used was ∼1 cm2 . The alloys contained from 11 at.% up to 46 at.% Nb or Ta and were donated by Prof. K. Hashimoto (Tohoku Institute of Technology, Sendai, Japan). The alloys’ layer was ∼600 nm thick. The materials were produced in a special clean room, to reduce the probability of crystallization of the sputtered material. High purity argon was used after being further purified by removing oxygen, water and dust. Pressure in the sputtering chamber was 10−4 Torr. Details concerning the fabrication procedure are given elsewhere [4]. The amorphous structure of the samples was confirmed by X-ray diffraction. The samples, before the electrochemical experiments, were rinsed with methanol and subsequently with distilled water, and eventually dried in air. Optical and scanning electron microscopy were used to examine the specimens before and after the experiments. Borate buffer solutions with added chloride (from 0 up to 1 M NaCl), prepared from analytical grade reagents, were used for the electrochemical measurements. The electrolytes were de-aerated with pure argon. Potentiostatic and potentiodynamic (dE/dt = 1 V h−1 ) experiments were performed to compare the anodic behavior of various Al–Ta and Al–Nb alloys. The potential was measured versus SCE (Saturated was Calomel Electrode). The apparent pitting potential, Enp estimated from the anodic polarization curves as shown in may be some 100 mV [17]. We are aware, however, that Enp
349
more positive than the true pitting potential [14,19], as discussed elsewhere[12–14,19]. Further, in order to assist understanding of the anodic behavior in the borate buffer electrolyte, samples of the alloys were polarized galvanostatically. The resultant films were thicker than the typical passive films formed during potentiostatic/potentiodynamic polarization, thus enabling detailed examination of the films by AES combined with depth profiling (Ar+ ions sputtering). Anodization was carried out at a constant c.d. of 1 mA cm−2 up to various voltages (up to 70 V) in a borate buffer electrolyte. Under these conditions, anodization of Al and its alloys proceeds with a high Faradic efficiency, in contrast to potentiostatic/potentiodynamic anodization where the efficiency is lower due to the contribution of the simultaneous anodic metal dissolution process [22]. The anodized samples were examined with AES to determine composition and estimate the thickness of the anodically formed film covering their surface. Composition profiles of surface films on Al–Ta/Nb alloys were measured with the aid of an AES 500 Auger electron spectrometer (RIBER) (equipped with an CI-50RB (RIBER) Ar+ ion sputtering gun) at the following parameters: vacuum < 5 × 10−10 Torr, primary energy 2 keV, modulation voltage 2 V, electron current ∼1 A at a beam spot of about φ 30 m. Depth profiling was carried out with a 4.5 kV Ar+ ion beam. Discontinuous sputtering was employed. In addition, a Microlab 3501 (Thermo VG Scientific) was used to determine the composition of the anodic surface layer and chemical state of the components, utilizing the highenergy resolution Auger electron spectroscopy function of the Microlab (energy resolution changeable from 0.6% down to 0.06%).
3. Results and discussion 3.1. Electrochemical measurements In our previous work [17], we have shown that the breakdown potential in 0.1 M NaCl for all the Al-based alloys is about 1 V higher than that for pure crystalline Al. The present results (Fig. 1) confirm that the beneficial effect of refractory metals is not confined only to a solution containing 0.1 M NaCl, but extends over a wide concentration range from 0.01 M up to 1 M NaCl. In Cl− -free borate buffer, we have shown that during potentiodynamic polarization the anodic c.d. is low (few A cm−2 ) starting from open circuit potential up to 2 V [17]. No c.d. rise was observed suggesting that an insulating film was formed at the surface, thus inhibiting the oxygen evolution reaction [17]. The present results confirm that sugges1 At the Physical Chemistry of Materials Center of the Institute of Physical Chemistry PAS and of the Faculty of Materials Science WUT, Warsaw, Poland.
350
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353
Fig. 1. Effect of chloride (Cl− ) concentration on pit nucleation potential ) for Al-based amorphous alloys, in borate buffer with NaCl content (Enp from 0.01 M up to 1 M NaCl.
tion. During galvanostatic anodization, films were formed supporting voltages as high as 70 V or more. Voltage grew linearly with time (see Fig. 2), suggesting that the anodic film grows at a constant rate, ∼100% efficiency, similar to the behavior of pure crystalline Al [8,20,21], as well as Al–Mo and Al–W AAs [8,14,18] in borate buffer. From the voltage versus time plots (Fig. 2) one can estimate the amount of charge Q(V) necessary to grow a film supporting a given voltage V, depending upon the composition of the alloy. An example of such an estimation for Q50 V versus at.% of a refractory metal is given in Fig. 3. One can see that the results for Al–Mo, Al–W and Al–Ta are very similar. The corresponding Q values for Al–Nb tend to be higher than that for the other AAs. Thus, more charge is necessary to produce an insulating film on Al–Nb AAs, than on the other AAs. The amount of anodic charge Q grows linearly with the concentration of the refractory metal (Fig. 3) except for Al–Ta
Fig. 2. Voltage vs. time plots for anodization of Al–Nb at a constant c.d. up to 25 V and 50 V.
Fig. 3. Effect of refractory metal on the amount of anodic charge necessary to produce a film on the surface of Al-based AAs, supporting voltage V = 50 V.
alloys, where a shallow local minimum at 26 at.% Ta is visible. The effect of refractory metal is generally small. This may suggest that the refractory metal plays a negligible role in the build-up of an anodic insulating layer. 3.2. Characterization of the anodic oxide film by Auger electron spectroscopy For a better understanding of the anodic behavior of Al-based AAs, surface analytical measurements were performed. Fig. 4 presents a typical composition profile for Al54 Nb46 after anodization up to 25 V for all layer components (a) and for metallic components only, with the exclusion of oxygen (b). Based on the individual AES spectra only two distinct domains can be distinguished within the profile: (I) an anodic oxide film containing Nb, Alox and O, and (II) the substrate containing Alm and Nb only. Apparently, both the components became oxidized here. They both took part in the anodic oxide film formation, thus migrating forwards as Al+3 and Nb+5 , whereas O−2 was migrating backwards. The thick arrows A, B, C, and D in Fig. 4 mark regions in which more detailed examination of the chemical state of the components has been performed with the aid of the Microlab 350. This instrument utilizes a spherical sector analyzer, with a kinetic energy resolution of ≤0.6%, as compared to the AES 500’s resolution of a few percent only. These measurements confirmed that Ta and Nb within the film were oxidized. Some examples of corresponding spectra for Al and Nb are given in Fig. 5a and b. One can see from these spectra (Fig. 5a) that within the regions A, B and C Al is oxidized, and there is no change in the spectrum in the course of sputtering. This suggests, in agreement with the literature [25], that there is no preferential sputtering of oxygen bound to Al. On the contrary, the spectra shown in Fig. 5b for Nb, suggest that Nb oxide may get decomposed in the course of depth sputtering,
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353
Fig. 4. Composition profiles of Al54 Nb46 AAs anodically oxidized up to 25 V (AES 500 Auger electron spectrometer-RIBER) (a) for all elements present in the films (b) with exclusion of oxygen (metallic components only). Two regions are distinguishable: (I) anodic oxide film, (II) the substrate. Thick arrows A, B, C, and D mark regions in which more detailed examination of the chemical state of the components has been performed with the aid of a Microlab 350 instrument.
as the spectra in the domains A, B and C differ from one another. The spectra in the reference region D are close to the spectra of metallic Al and Nb, as can be seen in Fig. 5a and b. Fig. 6 shows an example of the typical composition profiles for Al46 W54 AAs anodized at a constant c.d. up to 50 V; the data in Fig. 6b correspond to the depth distribution of the metallic components only. These results are not analogous to those presented in Fig. 4b. There is no refractory metal in region I. The relative concentration of refractory metal is higher in region II than in the substrate (region III). The suggestion can be made that various refractory metals contribute to anodically formed oxide layer in a dissimilar way. Mo and W do not undergo oxidation during a fast galvanostatic anodization process. Therefore, they remain accumulated within the inner part of the film (region II), while Al+3 migrates outwards to form region I. Hence, the relative concentration of Al in region II is lower than in region III. As the refractory metals Mo or W (which under the conditions investigated do not undergo ionisation) do not migrate in the high electric field they thus serve as natural markers of the primary surface of the sample.
351
Fig. 5. Examples of Auger spectra for Al54 Nb46 amorphous alloy covered by an anodic oxide film formed in borate buffer at I = 1 mA cm−2 up to 25 V (Microlab 350) (a) aluminum signal (KLL), (b) niobium signal (LMM). The spectra correspond to the regions (I) anodic oxide film (A, B and C) and (II) substrate (D), as indicated (compare the profile in Fig. 4. Reference spectra for pure Al (KLL) and pure Nb (LMM) are also given.
A lack of the refractory metals within the upper part of the oxide film (region I) and its relatively higher concentration within the inner part of the film (region II), found in the depth composition profiles, suggest that the enhanced stability of the Al–W and Al–Mo AAs may be due to their action as metal dissolution moderators as suggested by Marcus and not as a result of participation of refractory metals in the surface oxide film formation (see Refs. [14] and [18]). The chemical composition of the anodic films on Al–Nb and Al–Ta AAs is less complex, as already discussed. Apparently the high affinities of Nb and Ta to oxygen decide their participation in the anodic oxide film formation as passivity promoters [23]. The sputtering rate used for the samples under investigation needs comment. For Al–W AAs samples the sputtering rate used was two times faster than that used for Al–Ta and Al–Nb. The time of sputtering an anodic film for both samples is almost the same (compare Figs. 4 and 6). We have to take into account; however, that according to the literature, [25] the insulating properties of Al2 O3 formed anodically on Al can be characterized by a parameter 1.2 nm V−1 (i.e. the
352
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353
by a factor of five. Distinctly less charge is necessary to produce an anodic film supporting a given voltage on Al–Mo or Al–W than on Al–Nb or Al–Ta AAs. (3) Anodization of the Al–Mo or Al–W amorphous alloys containing Mo, W results in the formation of a bi-layer film. Its upper part is alumina which is an insulator. The lower part consists of a “mixture” of alumina and the refractory metal. This part is probably responsible for the improved stability of the passive state (high Enp), acting as a corrosion barrier. • Enhanced stability of these AAs may be due to the action of refractory metals as metal dissolution moderators rather, as suggested by Ph. Marcus but is not a result of participation of a refractory metal in the surface oxide film formation. (4) Anodization of Al–Nb or Al–Ta results in the formation of mixture of oxides of both constituent metals. Apparently, Nb and Ta act here as passivity promoters, in spite of their high metal–metal bond energy. The factor responsible for such a behaviour is probably their high affinity to oxygen, exceeding that of Al metal (see Table 1). Fig. 6. Composition profiles of anodically oxidized Al46 W54 AAs [18] (a) for all elements present in the films (b) with exclusion of oxygen (metallic components only). Three regions are distinguishable: (I) the upper part of the anodic film (no refractory metal is detectable there), (II) the inner part of the anodic film (relative W concentration is higher than within the substrate III), (III) the substrate.
thickness of the insulating film necessary to support a voltage of V = 1 V is 1.2 nm), whereas rather higher values up to 1.7 nm V−1 were found for anodic films formed on Al–Mo and Al–W AAs. Thus, we can reasonably estimate that the thickness of the anodic film on Al–W alloy (Fig. 6) was: 1.6 nm V−1 × 50 V = 80 nm, whereas on Al–Nb (Fig. 4) was only 1.6 nm V−1 × 25 V = 40 nm.
Acknowledgements The authors are grateful to: Prof. K. Hashimoto for donation of two series of highly homogeneous Al–Ta and Al–Nb AAs, and to the Alexander von Humboldt Foundation for donation of the AES 500 Spectrometer (RIBER) equipped with a Kernel programme for data acquisition and processing (for M.J.C.). Thanks are also due to Prof. A. Szummer for helpful discussions. This work was financially supported by grant KBN 4T08A 030 25, and by the Institute of Physical Chemistry PAS.
References 4. Summary and conclusions (1) Electrochemical measurements have shown that stability of the passive state of Al-based AAs in aggressive solutions containing Cl− is enhanced. The alloys undergo pitting corrosion at and above the breakdown potential, , which is over 1 V higher than that characteristic of Enp crystalline Al. The effect of Ta on Enp compares with that of Mo, while the effect of Nb is close to that of W. The extension of the stable passivity region is wider for Al–Nb and Al–W than for Al–Ta or Al–Mo in a wide Cl− concentration range from 0.01 M Cl− up to1 M Cl− . (2) Galvanostatic measurements in Cl− -free electrolyte have shown that an anodic film thus formed are insulators supporting voltages as high as 70 V or more. The amount of charge necessary to produce anodic film supporting a given voltage increases by only ∼30% when the concentration of the refractory metal increases
[1] M. Naka, K. Hashimoto, T. Masumoto, Corrosion 32 (1976) 146. [2] K.E. Heusler, D. Huerta, in: H. Matyja, P.G. Zielinski (Eds.), Proceedings of Summer School on Amorphous Metals, World Scientific, 1986, p. 302. [3] R. Zallen, The Physic of Amorphous Solids, Wiley, New York, 1983. [4] H. Yosioka, H. Habazaki, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci. 32 (1991) 313–327. [5] T.B. Massalski, H. Okamoto, P.R. Subramanian, L. Kacprzak (Eds.), Binary Alloy Phase Diagrams, second ed., ASM International. [6] J.B. Lumdsen, in: R.P. Frankenthal, J. Kruger (Eds.), Passivity of Metals, The Electrochemical Society Monograph, NJ, Pennigton, 1978, p. 730. [7] Ph. Marcus, Corros. Sci. 36 (1988) 2155. [8] H. Habazaki, M.A. Peaz, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, Corros. Sci. 38 (1999) 2053. [9] X. Zhou, G.E. Thompson, P. Skeldon, G.C. Wood, K. Shimizu, H. Habazaki, Corros. Sci. 41 (1999) 1599. [10] X. Zhou, H. Habazaki, G.E. Thompson, G.C. Wood, Corros. Sci. 38 (1996) 1563. [11] X. Zhou, G.E. Thompson, H. Habazaki, K. Shimizu, P. Skeldon, G.C. Wood, Thin Solid Films 293 (1997) 327.
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353 [12] M. Janik-Czachor, A. Wolowik, Z. Werner, Corros. Sci. 36 (1994) 1921. [13] A. Wolowik, M. Janik-Czachor, A. Szummer, in: P.M. Natishan, H.S. Isaacs, M. Janik-Czachor, V.A. Macagno, P. Marcus, M. Seo (Eds.), Proceedings of the Symposium on Passivity and its Breakdown, The Electrochemical Society, NJ, vol. 97 (26) (1998) 533. [14] A. Wolowik, M. Janik-Czachor, Mater. Sci. Eng., A 267 (1999) 301. [15] E. Sikora, B.A. Shaw, T. Miller, in: P.M. Natishan, H.S. Isaacs, M. Janik-Czachor, V.A. Macagno, P. Marcus, M. Seo (Eds.), Proceedings of the Symposium on Passivity and its Breakdown, The Electrochemical Society, NJ, vol. 97 (26) (1998) 654. [16] B. Shaw, W.C. Moshier, R. Went, P. Miller, E. Principe, 5th International Symposium on Electrochemical Methods in Corrosion Research EMCR’94, Lisbona, 1994. [17] M. Janik-Czachor, A. Jaskiewicz, P. Kedzierzawski, Z. Werner, Mater. Sci. Eng., A 358 (2003) 171.
353
[18] A. Wolowik, Thesis, Polish Academy of Science, Warsaw, 2001. [19] M. Janik-Czachor, J. Electrochem. Soc. 128 (1981) 513C. [20] H. Habazaki, P. Skeldon, G.E. Thompson, G.C. Wood, Philos. Mag. B 71 (1995) 81. [21] H. Habazaki, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, Corros. Sci. 37 (1997) 731. [22] T. Valand, K.E. Heusler, J. Electroanal. Chem. 149 (1983) 71. [23] Z. Werner, A. Jaskiewicz, M. Pisarek, M. Janik-Czachor, in preparation. [24] R. Kirchheim, B. Heine, S. Hofmann, H. Hofsass, Corros. Sci. 31 (1990) 573; R. Kirchheim, B. Heine, S. Hofmann, H. Hofsass, Corros. Sci. 29 (1989) 899. [25] D. Briggs, M.P. Seah, Practical Surface Analysis, vol. 1, second ed., John Wiley & Sons, 1990, p. 170.
Passivity and its breakdown in Al-based amorphous alloys M. Janik-Czachor a,∗ , A. Jaskiewicz a,b , M. Dolata a,c , Z. Werner a,d a
Physical Chemistry of Materials Center, Institute of Physical Chemistry, Polish Academy of Sciences, 01-224 Warsaw, Kasprzaka str. 44/52, Poland b Faculty of Materials Science, Warsaw University of Technology, Warsaw, Poland c Warsaw Agriculture University, Warsaw, Poland d The Andrzej Soltan Institute for Nuclear Studies, Swierk, Poland Received 8 September 2004; received in revised form 29 October 2004; accepted 15 November 2004
Abstract The effect of Cl− ions (0–1 M NaCl) on the anodic behaviour of two series of sputter-deposited Al–Ta (11–46 at.%) and Al–Nb (13–46 at.%) amorphous alloys was investigated in a neutral electrolyte by using electrochemical methods and the results were compared with those for Al–Mo and Al–W alloys. In solutions containing Cl− the alloys broke down at and above the pit nucleation potential Enp . They exhibited a stable passivity within a much wider (E > 1 V) potential range than crystalline Al. The anodically formed oxide films in Cl− -free electrolyte were characterized by surface analytical measurements including high-energy resolution AES. Galvanostatic polarization was used to produce thicker anodic films than those formed under natural conditions, to facilitate detailed examination of their depth composition profiles. It was found that the anodic films grown on Al–Ta and Al–Nb alloys had a homogeneous structure in contrast to those on Al–Mo and Al–W. Transition metal Ta or Nb was detected at any depth of the composition profile. A relative enrichment of refractory metal within the inner part of the anodic film was not found, for these alloys. Tentative mechanisms of enhanced stability of the passive state of the Al-based amorphous alloys (AAs) are discussed. © 2005 Elsevier B.V. All rights reserved. Keywords: Amorphous alloys; Al-based alloys; Passivity breakdown; Pit nucleation potential; Insulating films; High-energy resolution AES analysis
1. Introduction Amorphous alloys (AAs) offer a new possibility of studying the effects of alloying elements on corrosion and passivity in a wide concentration range. Their great compositional flexibility, makes it possible to study the chemical effects without any interaction of structural effect [1–4]. One can, for example, produce homogeneous Al-based AAs containing from ∼11 at.% up to 50 at.% of a refractory metal (Mo, W, Ta, or Nb) [4], while in the crystalline state the solubility of these refractory metals is below 1 at.% [5]. Refractory metals are known for their beneficial effect on the passivity of many alloys, both crystalline and amorphous, although the mechanism of their action is still a subject of discussion [6–19]. AAs thus offer the opportunity of gaining an insight into the role of these elements, when studying ∗
Corresponding author. Tel.: +48 22 3433225; fax: +48 22 6325276. E-mail address: [email protected] (M. Janik-Czachor).
0254-0584/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2004.11.017
their effect also at higher concentrations. In particular, one can check whether the preferential dissolution of Al and/or different transport rates within the passive film of the components determines the anodic behavior [24], or whether the effect of refractory metal is confined to the action of “metal dissolution moderators” rather than to “passivity promoters”, as suggested by Marcus [7]. As the previous results [14,18] strongly suggest the former effect for Al–W and Al–Mo, it was thus crucial to study Al–Ta and Al–Nb AAs in order to prove the generality of the “metal dissolution moderators” model. While the elements considered, Mo, W, Ta, and Nb, belong to a group of refractory metals with a high energy of Me–Me bond [7] (the Me–Me bond energy for all the mentioned refractory metals is between 170 and 220 kJ mol−1 , while the corresponding value for Al is <60 kJ mol−1 [7]), their affinity to oxygen differs considerably from one another. Table 1 lists the corresponding values of free enthalpy of formation of an oxide (G◦ ). The G◦ value for Al2 O3 is also given.
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353 Table 1 Gibbs free enthalpy of formation of relevant oxide Oxide
G◦ (kJ mol−1 )
MoO3 WO3 Al2 O3 Nb2 O5 Ta2 O5
−667.5 −764.2 −1582.72 −1765.86 −1910.99
While G◦ for molybdenum and tungsten oxide is lower than that for Al2 O3 , the corresponding values for Ta2 O5 and Nb2 O5 are much higher. Thus, the anodic behavior of the Ta and Nb AAs may be different than that of the other refractory metals mentioned above. To carefully investigate the depth distribution of both components within the protective surface film, thick anodic films formed on Al–Mo and Al–W AAs were studied [13,14,18] as model systems for typical thin passivating films. The same procedure was adopted in this work to study anodic films on Al–Ta and Al–Nb AAs, to provide a proper basis for comparison. The aim of this work was to investigate the stability of the passive state of Al–Ta and Al–Nb AAs, and to compare the results with those for Al–Mo/W AAs in order to gain insight into the beneficial effect of the refractory metals.
2. Experimental The materials used in the investigations were Al–Ta and Al–Nb amorphous alloys obtained by sputter deposition onto Pyrex glass substrates. The exposed surface area used was ∼1 cm2 . The alloys contained from 11 at.% up to 46 at.% Nb or Ta and were donated by Prof. K. Hashimoto (Tohoku Institute of Technology, Sendai, Japan). The alloys’ layer was ∼600 nm thick. The materials were produced in a special clean room, to reduce the probability of crystallization of the sputtered material. High purity argon was used after being further purified by removing oxygen, water and dust. Pressure in the sputtering chamber was 10−4 Torr. Details concerning the fabrication procedure are given elsewhere [4]. The amorphous structure of the samples was confirmed by X-ray diffraction. The samples, before the electrochemical experiments, were rinsed with methanol and subsequently with distilled water, and eventually dried in air. Optical and scanning electron microscopy were used to examine the specimens before and after the experiments. Borate buffer solutions with added chloride (from 0 up to 1 M NaCl), prepared from analytical grade reagents, were used for the electrochemical measurements. The electrolytes were de-aerated with pure argon. Potentiostatic and potentiodynamic (dE/dt = 1 V h−1 ) experiments were performed to compare the anodic behavior of various Al–Ta and Al–Nb alloys. The potential was measured versus SCE (Saturated was Calomel Electrode). The apparent pitting potential, Enp estimated from the anodic polarization curves as shown in may be some 100 mV [17]. We are aware, however, that Enp
349
more positive than the true pitting potential [14,19], as discussed elsewhere[12–14,19]. Further, in order to assist understanding of the anodic behavior in the borate buffer electrolyte, samples of the alloys were polarized galvanostatically. The resultant films were thicker than the typical passive films formed during potentiostatic/potentiodynamic polarization, thus enabling detailed examination of the films by AES combined with depth profiling (Ar+ ions sputtering). Anodization was carried out at a constant c.d. of 1 mA cm−2 up to various voltages (up to 70 V) in a borate buffer electrolyte. Under these conditions, anodization of Al and its alloys proceeds with a high Faradic efficiency, in contrast to potentiostatic/potentiodynamic anodization where the efficiency is lower due to the contribution of the simultaneous anodic metal dissolution process [22]. The anodized samples were examined with AES to determine composition and estimate the thickness of the anodically formed film covering their surface. Composition profiles of surface films on Al–Ta/Nb alloys were measured with the aid of an AES 500 Auger electron spectrometer (RIBER) (equipped with an CI-50RB (RIBER) Ar+ ion sputtering gun) at the following parameters: vacuum < 5 × 10−10 Torr, primary energy 2 keV, modulation voltage 2 V, electron current ∼1 A at a beam spot of about φ 30 m. Depth profiling was carried out with a 4.5 kV Ar+ ion beam. Discontinuous sputtering was employed. In addition, a Microlab 3501 (Thermo VG Scientific) was used to determine the composition of the anodic surface layer and chemical state of the components, utilizing the highenergy resolution Auger electron spectroscopy function of the Microlab (energy resolution changeable from 0.6% down to 0.06%).
3. Results and discussion 3.1. Electrochemical measurements In our previous work [17], we have shown that the breakdown potential in 0.1 M NaCl for all the Al-based alloys is about 1 V higher than that for pure crystalline Al. The present results (Fig. 1) confirm that the beneficial effect of refractory metals is not confined only to a solution containing 0.1 M NaCl, but extends over a wide concentration range from 0.01 M up to 1 M NaCl. In Cl− -free borate buffer, we have shown that during potentiodynamic polarization the anodic c.d. is low (few A cm−2 ) starting from open circuit potential up to 2 V [17]. No c.d. rise was observed suggesting that an insulating film was formed at the surface, thus inhibiting the oxygen evolution reaction [17]. The present results confirm that sugges1 At the Physical Chemistry of Materials Center of the Institute of Physical Chemistry PAS and of the Faculty of Materials Science WUT, Warsaw, Poland.
350
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353
Fig. 1. Effect of chloride (Cl− ) concentration on pit nucleation potential ) for Al-based amorphous alloys, in borate buffer with NaCl content (Enp from 0.01 M up to 1 M NaCl.
tion. During galvanostatic anodization, films were formed supporting voltages as high as 70 V or more. Voltage grew linearly with time (see Fig. 2), suggesting that the anodic film grows at a constant rate, ∼100% efficiency, similar to the behavior of pure crystalline Al [8,20,21], as well as Al–Mo and Al–W AAs [8,14,18] in borate buffer. From the voltage versus time plots (Fig. 2) one can estimate the amount of charge Q(V) necessary to grow a film supporting a given voltage V, depending upon the composition of the alloy. An example of such an estimation for Q50 V versus at.% of a refractory metal is given in Fig. 3. One can see that the results for Al–Mo, Al–W and Al–Ta are very similar. The corresponding Q values for Al–Nb tend to be higher than that for the other AAs. Thus, more charge is necessary to produce an insulating film on Al–Nb AAs, than on the other AAs. The amount of anodic charge Q grows linearly with the concentration of the refractory metal (Fig. 3) except for Al–Ta
Fig. 2. Voltage vs. time plots for anodization of Al–Nb at a constant c.d. up to 25 V and 50 V.
Fig. 3. Effect of refractory metal on the amount of anodic charge necessary to produce a film on the surface of Al-based AAs, supporting voltage V = 50 V.
alloys, where a shallow local minimum at 26 at.% Ta is visible. The effect of refractory metal is generally small. This may suggest that the refractory metal plays a negligible role in the build-up of an anodic insulating layer. 3.2. Characterization of the anodic oxide film by Auger electron spectroscopy For a better understanding of the anodic behavior of Al-based AAs, surface analytical measurements were performed. Fig. 4 presents a typical composition profile for Al54 Nb46 after anodization up to 25 V for all layer components (a) and for metallic components only, with the exclusion of oxygen (b). Based on the individual AES spectra only two distinct domains can be distinguished within the profile: (I) an anodic oxide film containing Nb, Alox and O, and (II) the substrate containing Alm and Nb only. Apparently, both the components became oxidized here. They both took part in the anodic oxide film formation, thus migrating forwards as Al+3 and Nb+5 , whereas O−2 was migrating backwards. The thick arrows A, B, C, and D in Fig. 4 mark regions in which more detailed examination of the chemical state of the components has been performed with the aid of the Microlab 350. This instrument utilizes a spherical sector analyzer, with a kinetic energy resolution of ≤0.6%, as compared to the AES 500’s resolution of a few percent only. These measurements confirmed that Ta and Nb within the film were oxidized. Some examples of corresponding spectra for Al and Nb are given in Fig. 5a and b. One can see from these spectra (Fig. 5a) that within the regions A, B and C Al is oxidized, and there is no change in the spectrum in the course of sputtering. This suggests, in agreement with the literature [25], that there is no preferential sputtering of oxygen bound to Al. On the contrary, the spectra shown in Fig. 5b for Nb, suggest that Nb oxide may get decomposed in the course of depth sputtering,
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353
Fig. 4. Composition profiles of Al54 Nb46 AAs anodically oxidized up to 25 V (AES 500 Auger electron spectrometer-RIBER) (a) for all elements present in the films (b) with exclusion of oxygen (metallic components only). Two regions are distinguishable: (I) anodic oxide film, (II) the substrate. Thick arrows A, B, C, and D mark regions in which more detailed examination of the chemical state of the components has been performed with the aid of a Microlab 350 instrument.
as the spectra in the domains A, B and C differ from one another. The spectra in the reference region D are close to the spectra of metallic Al and Nb, as can be seen in Fig. 5a and b. Fig. 6 shows an example of the typical composition profiles for Al46 W54 AAs anodized at a constant c.d. up to 50 V; the data in Fig. 6b correspond to the depth distribution of the metallic components only. These results are not analogous to those presented in Fig. 4b. There is no refractory metal in region I. The relative concentration of refractory metal is higher in region II than in the substrate (region III). The suggestion can be made that various refractory metals contribute to anodically formed oxide layer in a dissimilar way. Mo and W do not undergo oxidation during a fast galvanostatic anodization process. Therefore, they remain accumulated within the inner part of the film (region II), while Al+3 migrates outwards to form region I. Hence, the relative concentration of Al in region II is lower than in region III. As the refractory metals Mo or W (which under the conditions investigated do not undergo ionisation) do not migrate in the high electric field they thus serve as natural markers of the primary surface of the sample.
351
Fig. 5. Examples of Auger spectra for Al54 Nb46 amorphous alloy covered by an anodic oxide film formed in borate buffer at I = 1 mA cm−2 up to 25 V (Microlab 350) (a) aluminum signal (KLL), (b) niobium signal (LMM). The spectra correspond to the regions (I) anodic oxide film (A, B and C) and (II) substrate (D), as indicated (compare the profile in Fig. 4. Reference spectra for pure Al (KLL) and pure Nb (LMM) are also given.
A lack of the refractory metals within the upper part of the oxide film (region I) and its relatively higher concentration within the inner part of the film (region II), found in the depth composition profiles, suggest that the enhanced stability of the Al–W and Al–Mo AAs may be due to their action as metal dissolution moderators as suggested by Marcus and not as a result of participation of refractory metals in the surface oxide film formation (see Refs. [14] and [18]). The chemical composition of the anodic films on Al–Nb and Al–Ta AAs is less complex, as already discussed. Apparently the high affinities of Nb and Ta to oxygen decide their participation in the anodic oxide film formation as passivity promoters [23]. The sputtering rate used for the samples under investigation needs comment. For Al–W AAs samples the sputtering rate used was two times faster than that used for Al–Ta and Al–Nb. The time of sputtering an anodic film for both samples is almost the same (compare Figs. 4 and 6). We have to take into account; however, that according to the literature, [25] the insulating properties of Al2 O3 formed anodically on Al can be characterized by a parameter 1.2 nm V−1 (i.e. the
352
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353
by a factor of five. Distinctly less charge is necessary to produce an anodic film supporting a given voltage on Al–Mo or Al–W than on Al–Nb or Al–Ta AAs. (3) Anodization of the Al–Mo or Al–W amorphous alloys containing Mo, W results in the formation of a bi-layer film. Its upper part is alumina which is an insulator. The lower part consists of a “mixture” of alumina and the refractory metal. This part is probably responsible for the improved stability of the passive state (high Enp), acting as a corrosion barrier. • Enhanced stability of these AAs may be due to the action of refractory metals as metal dissolution moderators rather, as suggested by Ph. Marcus but is not a result of participation of a refractory metal in the surface oxide film formation. (4) Anodization of Al–Nb or Al–Ta results in the formation of mixture of oxides of both constituent metals. Apparently, Nb and Ta act here as passivity promoters, in spite of their high metal–metal bond energy. The factor responsible for such a behaviour is probably their high affinity to oxygen, exceeding that of Al metal (see Table 1). Fig. 6. Composition profiles of anodically oxidized Al46 W54 AAs [18] (a) for all elements present in the films (b) with exclusion of oxygen (metallic components only). Three regions are distinguishable: (I) the upper part of the anodic film (no refractory metal is detectable there), (II) the inner part of the anodic film (relative W concentration is higher than within the substrate III), (III) the substrate.
thickness of the insulating film necessary to support a voltage of V = 1 V is 1.2 nm), whereas rather higher values up to 1.7 nm V−1 were found for anodic films formed on Al–Mo and Al–W AAs. Thus, we can reasonably estimate that the thickness of the anodic film on Al–W alloy (Fig. 6) was: 1.6 nm V−1 × 50 V = 80 nm, whereas on Al–Nb (Fig. 4) was only 1.6 nm V−1 × 25 V = 40 nm.
Acknowledgements The authors are grateful to: Prof. K. Hashimoto for donation of two series of highly homogeneous Al–Ta and Al–Nb AAs, and to the Alexander von Humboldt Foundation for donation of the AES 500 Spectrometer (RIBER) equipped with a Kernel programme for data acquisition and processing (for M.J.C.). Thanks are also due to Prof. A. Szummer for helpful discussions. This work was financially supported by grant KBN 4T08A 030 25, and by the Institute of Physical Chemistry PAS.
References 4. Summary and conclusions (1) Electrochemical measurements have shown that stability of the passive state of Al-based AAs in aggressive solutions containing Cl− is enhanced. The alloys undergo pitting corrosion at and above the breakdown potential, , which is over 1 V higher than that characteristic of Enp crystalline Al. The effect of Ta on Enp compares with that of Mo, while the effect of Nb is close to that of W. The extension of the stable passivity region is wider for Al–Nb and Al–W than for Al–Ta or Al–Mo in a wide Cl− concentration range from 0.01 M Cl− up to1 M Cl− . (2) Galvanostatic measurements in Cl− -free electrolyte have shown that an anodic film thus formed are insulators supporting voltages as high as 70 V or more. The amount of charge necessary to produce anodic film supporting a given voltage increases by only ∼30% when the concentration of the refractory metal increases
[1] M. Naka, K. Hashimoto, T. Masumoto, Corrosion 32 (1976) 146. [2] K.E. Heusler, D. Huerta, in: H. Matyja, P.G. Zielinski (Eds.), Proceedings of Summer School on Amorphous Metals, World Scientific, 1986, p. 302. [3] R. Zallen, The Physic of Amorphous Solids, Wiley, New York, 1983. [4] H. Yosioka, H. Habazaki, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci. 32 (1991) 313–327. [5] T.B. Massalski, H. Okamoto, P.R. Subramanian, L. Kacprzak (Eds.), Binary Alloy Phase Diagrams, second ed., ASM International. [6] J.B. Lumdsen, in: R.P. Frankenthal, J. Kruger (Eds.), Passivity of Metals, The Electrochemical Society Monograph, NJ, Pennigton, 1978, p. 730. [7] Ph. Marcus, Corros. Sci. 36 (1988) 2155. [8] H. Habazaki, M.A. Peaz, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, Corros. Sci. 38 (1999) 2053. [9] X. Zhou, G.E. Thompson, P. Skeldon, G.C. Wood, K. Shimizu, H. Habazaki, Corros. Sci. 41 (1999) 1599. [10] X. Zhou, H. Habazaki, G.E. Thompson, G.C. Wood, Corros. Sci. 38 (1996) 1563. [11] X. Zhou, G.E. Thompson, H. Habazaki, K. Shimizu, P. Skeldon, G.C. Wood, Thin Solid Films 293 (1997) 327.
M. Janik-Czachor et al. / Materials Chemistry and Physics 92 (2005) 348–353 [12] M. Janik-Czachor, A. Wolowik, Z. Werner, Corros. Sci. 36 (1994) 1921. [13] A. Wolowik, M. Janik-Czachor, A. Szummer, in: P.M. Natishan, H.S. Isaacs, M. Janik-Czachor, V.A. Macagno, P. Marcus, M. Seo (Eds.), Proceedings of the Symposium on Passivity and its Breakdown, The Electrochemical Society, NJ, vol. 97 (26) (1998) 533. [14] A. Wolowik, M. Janik-Czachor, Mater. Sci. Eng., A 267 (1999) 301. [15] E. Sikora, B.A. Shaw, T. Miller, in: P.M. Natishan, H.S. Isaacs, M. Janik-Czachor, V.A. Macagno, P. Marcus, M. Seo (Eds.), Proceedings of the Symposium on Passivity and its Breakdown, The Electrochemical Society, NJ, vol. 97 (26) (1998) 654. [16] B. Shaw, W.C. Moshier, R. Went, P. Miller, E. Principe, 5th International Symposium on Electrochemical Methods in Corrosion Research EMCR’94, Lisbona, 1994. [17] M. Janik-Czachor, A. Jaskiewicz, P. Kedzierzawski, Z. Werner, Mater. Sci. Eng., A 358 (2003) 171.
353
[18] A. Wolowik, Thesis, Polish Academy of Science, Warsaw, 2001. [19] M. Janik-Czachor, J. Electrochem. Soc. 128 (1981) 513C. [20] H. Habazaki, P. Skeldon, G.E. Thompson, G.C. Wood, Philos. Mag. B 71 (1995) 81. [21] H. Habazaki, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, Corros. Sci. 37 (1997) 731. [22] T. Valand, K.E. Heusler, J. Electroanal. Chem. 149 (1983) 71. [23] Z. Werner, A. Jaskiewicz, M. Pisarek, M. Janik-Czachor, in preparation. [24] R. Kirchheim, B. Heine, S. Hofmann, H. Hofsass, Corros. Sci. 31 (1990) 573; R. Kirchheim, B. Heine, S. Hofmann, H. Hofsass, Corros. Sci. 29 (1989) 899. [25] D. Briggs, M.P. Seah, Practical Surface Analysis, vol. 1, second ed., John Wiley & Sons, 1990, p. 170.