Anodic behaviour of Fe73.5Si13.5-xAlxB9Nb3Cu1 (x = 0–2) amorphous, nanostructured and crystalline alloys

Scripta mater. 44 (2001) 507–512 www.elsevier.com/locate/scriptamat

ANODIC BEHAVIOUR OF Fe73.5Si13.5-xAlxB9Nb3Cu1 (X ⴝ 0 –2) AMORPHOUS, NANOSTRUCTURED AND CRYSTALLINE ALLOYS M.G. Alvarez U.A. Materiales, Centro Ato´mico Constituyentes, Comisio´n Nacional de Energı´a Ato´mica, Avda. Libertador 8250, 1429 Buenos Aires, Argentina

S.M. Vazquez, J. Moya and H. Sirkin Laboratoria de So´lidos Amorfos, Departamento de Fı´sica, Facultad de Ingenierı´a UBA, Paseo Colo´n 850, 1063 Buenos Aires, Argentina (Received February 2, 2000) (Accepted in revised form July 7, 2000) Keywords: Nanocrystalline; Metallic glasses; Corrosion; Electrochemical techniques Introduction The development of nanostructural alloys has led to a significant improvement of soft magnetic materials (1). One of the most prominent examples is nanocrystalline FeSiBNbCu alloys. Owing to the formation of nanocrystallites embedded in an amorphous matrix, the magnetic anisotropy is substantially reduced and the magnetostrictive effects are eliminated (2). The magnetic properties and structural changes of FeSiBCuNb alloys have been studied extensively (2– 6). Owing to the widespread commercial utilisation envisaged for nanostructured materials, also the corrosion performance of FeSiBCuNb nanocrystalline alloys must be evaluated. In the present work, a comparative study of the corrosion behaviour of Fe73.5Si13.5B9Cu1Nb3 amorphous, nanocrystalline and crystalline alloys in a chloride containing solution has been performed by means of electrochemical techniques. The influence of the microstructural changes caused by primary and secondary crystallisation on the polarization behaviour of the amorphous alloy was examined. The effect of the partial replacement of silicon by aluminium (1% and 2% Al) on crystallisation and corrosion behaviour has also been analyzed. Experimental Method Materials Amorphous alloys of nominal composition Fe73.5Si13.5B9Nb3Cu1, Fe73.5Si12.5Al1B9Nb3Cu1 and Fe73.5Si11.5Al2B9Nb3Cu1 nanocrystalline alloys specimens were obtained by the melt-spinning technique in the shape of ribbons 1.0 mm wide and 20 –30 ␮m thick. The nanocrystalline and crystalline alloy specimenswere prepared by annealing the amorphous ribbons for 1hr at 560 °C and 750 °C in a vacuum of 4 ⫻ 10⫺5 mbar. X-ray diffraction (XRD) measurements on as-quenched and annealed samples were carried out using Cu K␣ radiation. 1359-6462/01/$–see front matter. © 2001 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(00)00627-8

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Figure 1. XRD-patterns of Fe73.5Si13.5-xB9Nb3Cu1Alx with x ⫽ 0 and x ⫽ 2% after annealing at 560 and 750 °C.

The nanocrystalline state is characterised by an ultrafine grain structure of ␣-Fe(Si) with a grain size of approximately 10 –20 nm, embedded in a still amorphous minority matrix. The Fe-Si grains grow following a DO3 structure (bcc Fe3Si type). ␣-Fe(Si,Al) bcc structure was also detected in the aluminium containing alloys. The amorphous matrix phase mainly consists of Fe, Nb and B. Copper enriched particles were also detected in the intergranular regions. After annealing at 750 °C, the remaining amorphous phase crystallises and in addition to ␣-Fe(Si), other boron-containing crystalline phases were detected. Figure 1 shows the XRD patterns of Fe73.5Si13.5B9Nb3Cu1 and Fe73.5Si11.5Al2B9Nb3Cu1 alloys annealed at 560 and 750 °C. After annealing at 560 °C only ␣-Fe(Si) crystalline phase was identified. When the annealing temperature increase to 750 °C some new crystalline peaks attributed to Fe23B6 and Fe3B were observed. The Fe3B peaks did not match very well with bibliography data. This effect was probably due to the presence of Nb-containing phases, such as (FeNb)2B and (FeNb)3B, as reported by Herzer (7). Electrochemical Tests The ribbons were cut into pieces 3 cm in length. A copper wire lead was attached with silver paint to one end of the samples. A Pyrex glass tube 1.5 mm in diameter was used to isolate the copper wire and the joint. The samples were held in place by covering a portion of the tube in contact with the ribbon with an epoxy resin. Electrical continuity was assured by filling up the tube with mercury. The exposed area of the samples was ca. 0.5 cm2. Before each experiment the samples were degreased with ethyl acetate, washed with ethyl alcohol, and distilled water. Polarisation measurements were performed by the potentiodynamic method in 1M NaCl, pH ⫽ 9.0, solution at room temperature. A Pyrex glass cell with a platinum counterelectrode was employed. Potentials were measured through a Luggin capillary, with a saturated calomel electrode. All potentials are reported in the normal hydrogen electrode scale (NHE). The solutions were prepared with analytical grade reagents and double distilled water and were de-aerated before introducing the specimen and during the experiments by bubbling purified nitrogen. The pH of the NaCl solution was adjusted by addition of NaOH. Potentiodynamic polarisation curves were obtained at a scan rate of 0.2 mV/s using a PAR 173 potentiostat with a PAR 175 voltage scan generator. The current and potential were continuously recorded with a Houston 2000 X-Y recorder. The samples were allowed to reach a stationary open

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Figure 2. Potentiodynamic anodic polarization curves of Fe73.5Si13.5-xB9Nb3 AlxCu1 (x ⫽ 0 –2) amorphous alloys in NaCl 1.0 M, pH ⫽ 9, solution.

circuit potential (90 min). Then, a potential value 200 mV lower than the corrosion potential was applied for 5 minutes and the potentiodynamic scan was initiated. Scanning electron microscopy (SEM) was used to characterise the specimen before and after the polarisation tests. Results The shape of the polarisation curves of Fe73.5Si13.5B9Nb3Cu1, Fe73.5Si12.5Al1B9Nb3Cu1 and Fe73.5Si11.5Al2B9Nb3Cu1 amorphous alloys in 1M NaCl, pH ⫽ 9, solution was similar for all three alloys (Figure 2). The corrosion potential was found to be around ⫺0.58 VNHE. Upon increasing the potential above the corrosion potential, the current density increased monotonically with the electrode potential. At about 0.1 VNHE an anodic peak, A1, associated to an active-passive transition was observed. With further increase in the potential a passive region was found, where the current density was of the order of 10⫺4 A/cm2. A goldish film was observed on the surface of the specimens and no further current variation was detected although the potential was increased up to 2.0 VNHE . As for the high silicon-iron alloys, the passive region observed at potentials higher than 0.1 VNHE could be associated to the development of a corrosion resistant film containing a large proportion of silica (8). The beneficial effect of the addition of silicon on the corrosion resistance of iron base amorphous metal-metalloid alloys owing to the formation of a silicon-enriched surface film has already been reported by several authors (9 –11). In the present work, the formation of a silicon rich-film was confirmed by the microscopic observation of the samples after polarisation. A thick film was found to be present on the surface of the specimens. The EDS analysis indicated that silicon was the main component of the film. In Figures 3 and 4 potentiodynamic polarisation curves of the amorphous, nanocrystalline and crystalline configurations of the three alloys in 1M NaCl, pH ⫽ 9, solution are compared. Several differences could be observed between the anodic behaviour of the nanocrystalline and crystalline alloys and that of their amorphous counterpart. The polarisation curves of all three nanocrystalline alloys showed two anodic peaks, labelled A1 and A2, both associated to an active-passive transition. The first activation peak, A1, occurred at nearly the same potential as for the amorphous state (0.1 VNHE). In the potential range between A1 and A2 a pseudo-passive region was found, where the current density was of the order of 10⫺3 A/cm2. The second activation peak, A2, at a potential of about 0.32 VNHE, was very small. At potentials higher than A2, the current density attained a value coincident with the one measured for the amorphous alloys in the same potential range. The corrosion potential of the nanocrystalline specimens were less negative than those measured for the amorphous state. The increase in the corrosion potential was attributed to the increased local

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Figure 3. Potentiodynamic polarization curves of Fe73.5Si13.5B9Nb3Cu1 amorphous, nanocrystalline and crystalline alloys in NaCl 1N, pH ⫽ 9, solution.

galvanic cell activity between the ␣-Fe(Si) nanocrystals and the residual amorphous phase, which were Nb, B and Cu enriched. The presence of copper clusters in the intergranular amorphous phase will favour the cathodic reaction of hydrogen evolution, leading to a displacement of the corrosion potential towards higher potential values. The polarisation curves of all three crystalline alloys (Fig. 3– 4) exhibited two activation peaks, A⬘1 and A⬘2, at more negative and positive potentials, respectively, than the activation peak A1 detected by anodic polarisation in the amorphous and nanostructured materials. The second activation peak, A⬘2, was higher and broader than the peak A2 observed in the polarisation curves of the nanocrystalline alloys. The similarity in shape of the polarisation curves measured for all three alloys in the amorphous as well as in the nanocrystalline and crystalline states shows that the partial replacement of silicon by aluminium (up to 2% Al) has no substantial effect on the anodic behaviour of the alloys in NaCl solutions. The major change was a displacement of the corrosion potential of the nanocrystalline alloys towards the less noble direction with increasing Al content. Discussion The results obtained show that the microstructure of the material has a noticeable effect on the polarisation behaviour. The amorphous alloys have a homogeneous chemical composition that can supply silicon to build up a silicon-rich film over the entire alloy surface. The shape of the corresponding polarisation curves (Fig. 2) suggests that, as for the high silicon-iron alloys, some corrosion of the base metal must necessary take place before the silica film can form (8). A significant dissolution of the more active non-passivating elements occurred as the potential was increased above the corrosion

Figure 4. Potentiodynamic polarization curves of Fe73.5Si12.5B9Nb3Al1Cu1 (a) and Fe73.5Si11.5B9Nb3Al2Cu1 (b) amorphous, nanocrystalline and crystalline alloys in NaCl 1M, pH ⫽ 9, solution.

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potential, followed by a marked decrease in anodic activity resulting from the formation of a silicon rich film. Thus, it can be assumed that the activation peak A1 was due to the oxidation of iron from the amorphous alloys. Fe 3 Fe2⫹ ⫹ 2e⫺

(1)

After heat treatment at 560 °C, primary crystallisation of ␣-Fe(Si) nanocrystals occurred. The volume fraction of the nanoscale phase was about 70% (3,5,6), with a 19%–22% Si concentration. Since presumably almost all Si has partitioned into the crystalline phase, the Si concentration was significantly different from place to place. The results of the polarisation experiments showed that the heterogeneous distribution of silicon in the nanocrystalline configuration is deleterious to the development of a stable passivating film. The passive film formed on the alloy surface was dictated by the chemical nature of the most abundant ␣-Fe(Si) phase. The activation peak A1 was present in the polarisation curves of the amorphous alloys. Consequently, the peak A1 in the nanocrystalline configuration could be associated with the dissolution of iron from the ␣-Fe(Si) phase and the build up of a silicon rich film.

␣ -Fe(Si) 3 Fe2⫹ ⫹ 2e⫺

(2)

On the other hand, the local depletion of Si in the remaining amorphous grain boundary phase seemed to prevent a continuous build up of the silicon rich layer. Dissolution of iron from the silicon depleted amorphous matrix continues at potentials higher than A1. It was only at potentials above the second activation peak, A2, that a stable passive film could be formed on the alloy surface. Annealing at 750 °C leads to full crystallisation. In addition to the ␣-Fe(Si) phase the presence of various intermetallic compounds (Fe3B and Fe23B6) and presumably, (FeNb)2B and/or (FeNb)3B, produced by the transformation of the remaining amorphous phase were detected. The grain size of the ␣-Fe(Si) phase in the full crystallised state was higher than in the nanocrystalline state (⬃ 22 and 9 nm respectively calculated with Scherre formula (12)). ␣-Fe(Si,Al) phase was also detected in the aluminium-containing alloys. As mentioned above, the corresponding polarisation curves showed two activation peaks, A⬘1 and A⬘2, at more negative and positive potentials, respectively, than the activation peak A1 resulting from the dissolution of iron in the amorphous and nanocrystalline configurations. As in the amorphous and nanocrystalline alloys, the first peak, A⬘1 could be related to the oxidation of iron from the ␣-Fe(Si,Al) phase according to reaction (2). The second activation peak, A⬘2, was absent in the polarisation curves obtained for the amorphous alloys. Thus, A⬘2 in the crystalline configuration can be attributed to the dissolution of the secondary crystallisation products, Fe3B and/or Fe23B6, which can not be separated by the electrochemical signal. FexB 3 xFe2⫹ ⫹ 2x e⫺

(3)

The presence of a second activation peak in the anodic polarisation curves after devitrification owing to the formation of Fe2B/Fe3B precipitates was previously reported for Fe80B14Si6, Fe80B20 and Fe77.5B15Si7.5 alloys in borate buffer solutions of pH ⫽ 9.2 and pH ⫽ 8.7 (13,14). Two activation peaks were also observed but not discussed in the polarisation curve of crystalline Fe73.5Si13.5B9Nb3Cu1 alloy in buffered acetate solutions containing different amounts of chloride (15). The present results, concerning the effect of microstructural changes on the anodic behaviour of FeSiAlBCuNb in chloride solutions, did not allow quantification of the volume fraction of the individual phases. However, the different shape and size of the two activation peaks in the nanocrystalline and crystalline states was probably associated with differences in composition and volume fraction of the various crystalline phases formed on annealing at 750 °C.

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Conclusions ● ● ● ●

Amorphous FeSiAlBNbCu alloys passivate in 1M NaCl, pH ⫽ 9, solution owing to the formation of a silicon-enriched surface film. The heterogeneous microstructure formed by primary and secondary crystallisation has a deleterious effect on the stability of the passive film. The partial replacement of silicon by aluminium (at least up to 2% Al) has no significant effect on the polarisation behaviour of the alloys in the amorphous as well as devitrified states. A good correlation is found between the microstructural changes and the electrochemical response of FeSiAlBNbCu alloys. References

1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

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