On the key role of Cu on the oxidation behavior of Cu–Ni–Fe based anodes for Al electrolysis

Corrosion Science 101 (2015) 105–113

Contents lists available at ScienceDirect

Corrosion Science journal homepage: www.elsevier.com/locate/corsci

On the key role of Cu on the oxidation behavior of Cu–Ni–Fe based anodes for Al electrolysis E. Gavrilova a , G. Goupil a , B. Davis b , D. Guay a , L. Roué a,∗ a b

INRS-Énergie, Matériaux et Télécommunication, 1650 Blvd Lionel-Boulet, Varennes, Québec J3X 1S2, Canada Kingston Process Metallurgy Inc., 759 Progress Avenue, Kingston, Ontario K7M 6N6, Canada

a r t i c l e

i n f o

Article history: Received 9 June 2015 Received in revised form 2 September 2015 Accepted 6 September 2015 Available online 9 September 2015 Keywords: A. Alloy B. SEM B. X-ray diffraction B. Galvanostatic C. High temperature corrosion

a b s t r a c t In order to prepare anode materials for testing for Al electrolysis, monophased Cu–Ni–Fe alloys were produced by high-energy ball milling with different Cu contents (65, 55 and 45 wt%) and a Ni/Fe mass ratio fixed at 1.33. A Cu-free alloy (57Ni43Fe) was also evaluated for comparison. Their oxidation behavior was studied at 700 ◦ C under O2 atmosphere and in Al electrolysis conditions. Under both conditions, it is shown that the presence of CuO in the outer layer is critical for the formation of an inner protective layer of NiFe2 O4 and depends critically on the Cu content of the alloy. © 2015 Elsevier Ltd. All rights reserved.

1. Introduction “Inert” anodes have been developed for several years as an alternative to the consumable carbon anodes presently used in the Hall–Héroult electrolysis cells for primary Al production [1]. One of the key features of inert anodes is their production of O2 instead of CO2 during the Al electrolysis process, inducing a significant decrease of the greenhouse gas emissions. As inert anodes are not rapidly consumed during electrolysis in contrast to carbon anodes, a decrease of the Al production costs is also expected by preventing frequent anode lowering and replacement operations. In addition, their implementation will eliminate the carbon anode production plant and would result in smaller Al production plants because more compact vertical configuration cells could be used. Obtaining an inert anode satisfying all the requirements to be viable under the extreme Al electrolysis conditions is very challenging. Among the possible anode materials (metals, ceramics and cermets) [2,3], metallic anodes are the most relevant candidates because they offer high electrical conductivity, excellent thermal shock resistance, mechanical robustness, ease of manufacture and simplicity of electrical connection to the current lead. However, to date, no long-term viable metallic anodes have been found due to their insufficient corrosion–dissolution resistance in the traditional

∗ Corresponding author. E-mail address: [email protected] (L. Roué). http://dx.doi.org/10.1016/j.corsci.2015.09.006 0010-938X/© 2015 Elsevier Ltd. All rights reserved.

electrolytic bath (NaF–AlF3 based electrolyte at 960 ◦ C). As such, a decrease in the operating temperature is needed to reach an acceptable dissolution rate (<1 cm year−1 ) for metallic inert anodes. In this context, the use of a low temperature KF–AlF3 -based electrolyte displaying a high alumina solubility (∼5 wt% at 700 ◦ C) could offer new opportunities for metal anodes [4–6]. As shown by Beck et al. [7,8], Cu–Ni–Fe based alloys present promising performance in low-temperature (700–850 ◦ C) aluminasaturated cryolite. This is mainly attributed to their ability to form a low solubility NiFe2 O4 protective surface scale during Al electrolysis. However, as-cast Cu–Ni–Fe alloys present a bi-phased structure, in which the Fe-rich phase is preferentially attacked during Al electrolysis, leading to the creation of deleterious corrosion tunnels [9]. We have recently shown that highly homogeneous Cu–Ni–Fe alloys can be obtained over a large composition range by mechanical alloying [10]. Moreover, this homogeneity is retained during their cold-spray deposition to produce large area electrodes [11] consisting of a reusable anode core and a tailored surface alloy. Best results were obtained with the Cu65 Ni20 Fe15 and (Cu65 Ni20 Fe15 )98.6 O1.4 (wt%) materials, which exhibit good corrosion resistance in low temperature (700 ◦ C) KF–AlF3 -based electrolyte, resulting in the production of Al with a purity of 99.3 and 99.8%, respectively [12]. However, despite the formation of a NiFe2 O4 -rich surface layer, the outward diffusion of Cu in Cu oxides is not fully impeded, inducing the formation of a porous Cudepleted region at the oxide/alloy interface. As a result, electrolyte

106

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

Table 1 Composition of the Cu–Ni–Fe alloys determined by EDX. Sample

[Cu] (wt%) Nominal

Cu-65 Cu-55 Cu-45 Cu-0

65 55 45 0

[Ni] (wt%) Measured

64.3 54.2 43.8 0

Nominal 20 25.7 31.4 57

Table 2 Lattice parameter and crystallite size of the Cu–Ni–Fe alloys determined from Rietveld refinement of the XRD patterns.

[Fe] (wt%) Measured 19.6 24.7 29.4 54.5

Nominal 15 19.3 23.6 43

Measured 16.1 21.1 26.8 45.5

penetration occurs in the scale, which favors the progressive formation of an iron fluoride-rich layer at the oxide/alloy interface and FeF2 inclusions in the bulk anode [13]. This is assumed to be responsible for the progressive increase of the cell voltage during prolonged electrolysis [11]. The objective of the present study is to evaluate the impact of a decrease in the Cu content in Cu–Ni–Fe anodes on the deleterious electrolyte penetration in the oxide scale. It will be highlighted that Cu plays a key role in the formation of protective NiFe2 O4 -rich surface layer. 2. Experimental Three Cu–Ni–Fe alloys were prepared with different Cu contents (65, 55 and 45 wt%) and a Ni/Fe mass ratio fixed at 1.33, as indicated in Table 1. A Cu-free alloy Ni57 Fe43 was also evaluated for comparison. The alloys were synthesized from elemental Cu, Ni and Fe powders (Cu purity ≥99.5%, Ni and Fe purity ≥99.9, −325 mesh) by high-energy ball milling using a vibratory mill (SPEX 8000) as described in detail elsewhere [14]. The milling time was 10 h. The composition of the as-milled powders determined by energy dispersive X-ray (EDX) analysis was in accordance (within 1–2 wt%) with their nominal composition (Table 1). A slight Fe enrichment was observed due to the erosion of the milling tools, which tends to increase with the decreasing alloy Cu content. Their O content determined by LECO was 0.3–0.4 wt%. The milled powders were then consolidated by a softening + cold-pressing + sintering procedure as described elsewhere [14]. The resulting pellets had a diameter of 11.3 mm, a thickness

Sample

Cu-65

Cu-55

Cu-45

Cu-0

Lattice parameter (Å) Cristallite size (nm)

3.599 21

3.594 23

3.589 28

3.574 53

of ∼5 mm for the electrolysis tests or ∼1 mm for the oxidation tests and a porosity of 3–6%. The crystalline structure of the consolidated powders was determined by X-ray diffraction (XRD) using a Bruker D8 diffractometer with Cu K␣ radiation, and refined using FullProf software. The oxidation behavior of the consolidated samples was evaluated by thermogravimetric analysis (TGA) using a Thermax 500 equipment. The oxidation experiments were conducted for 20 h at 700 ◦ C under 1 atm 80% Ar −20% O2 with a flow rate of 240 cm3 min−1 . After oxidation, the oxide scale was analyzed by XRD. The samples were then enrobed in epoxy resin and polished to examine their cross section by SEM and EDX mapping. Electrolysis tests were performed at an anodic current density of 0.5 A cm−2 for 20 h in alumina-saturated KF–AlF3 electrolyte (50 wt% AlF3 + 45 wt% KF + 5 wt% Al2 O3 ) at 700 ◦ C. The electrochemical reactor is presented in detail elsewhere [15]. The electrolysis experiment was repeated at least twice for each anode composition, confirming the repeatability of the cell voltage evolution. After electrolysis, the anodes were examined by cross sectional SEM–EDX. The purity of the produced Al was determined from neutron activation analyses. 3. Results and discussion 3.1. Crystalline structure Fig. 1 shows the XRD patterns of the consolidated powders. For all samples, only one series of peak corresponding to a facecentered-cubic (fcc) phase (␥-phase) is observed, confirming that the alloying process is completed. This also suggests that no spinodal decomposition occurs in contrast to what is usually observed for as-cast Cu–Ni–Fe alloys, which segregate into  1 Cu-rich and ␥2 Fe–Ni-rich fcc phases during the cooling step [16,17]. This is also confirmed by the absence of chemical segregation and dendritic domains characteristic of a spinodal decomposition on the backscattered electron (BSE) images and EDX maps of the samples (not shown). As shown in Table 2, a diminution of the alloy lattice parameter is observed as its Cu content decreases, in accordance with the larger atomic radius of Cu (1.28 Å) compared to that of Ni (1.24 Å) and Fe (1.26 Å). An increase of the crystallite size from 21 to 53 nm is observed with the decreasing Cu content (Table 2), which may reflect some variation in the mechanical properties of the alloy with its composition affecting the powder fracturing process upon milling. 3.2. Oxidation behavior

Fig. 1. XRD patterns of the consolidated Cu-65, Cu-55, Cu-45 and Cu-0 powders.

Fig. 2A shows the TGA curves for the Cu-65, Cu-55, Cu-45 and Cu-0 samples at 700 ◦ C under 1 atm Ar −20%O2 . Fig. 2B displays the same data plotted as (weight gain)2 versus time. All the curves present a shape characteristic of a diffusion-controlled oxidation process. The oxidation rate decreases significantly as the Cu content in the Cu–Ni–Fe alloy decreases. The parabolic rate constant kp determined from Fig. 2B between 5 and 20 h of oxidation is about 2.8, 1.2 and 0.28 × 10−10 g2 cm−4 s−1 for the Cu-65, Cu-55, Cu-45 samples, respectively. These values are not very far from those obtained by Gallino et al. on various Cu–Ni–Fe alloys at 800 ◦ C under

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

-2

Weight gain (mg cm )

5

(A)

Cu-65

4

Cu-55

3

Cu-0

2

Cu-45 1 0

0

5

15

20

Cu-65

20

2

2

-4

10

Time (h)

(B) Weight gain (mg cm )

107

10

Cu-55 Cu-0 Cu-45

0

0

5

10

15

Fig. 3. XRD patterns of the oxidized (after TGA) Cu-65, Cu-55, Cu-45 and Cu-0 samples.

20

Time (h) Fig. 2. Weight gain (A) and square of weight gain (B) as a function of the oxidation time at 700 ◦ C under 1 atm Ar: O2 (80:20) for the Cu-65, Cu-55, Cu-45 and Cu-0 samples.

1 atm O2 [18]. The oxidation rate of the Ni–Fe alloy is higher than that of the Cu-45 alloy, with a kp value of 0.57 × 10−10 g2 cm−4 s−1 . As shown in Fig. 3, the XRD patterns of the post-TGA samples are similar for the three Cu–Ni–Fe alloys. The surface oxide scale is mainly constituted of CuO. Cu2 O and NiFe2 O4 phases are also detected. The peaks of the Cu2 O phase (at ∼36 and ∼42◦ ) become less intense with the decreasing Cu content in the alloy. The presence of NiO and Fe3 O4 phases in the oxide scale cannot be excluded since their characteristic peaks closely overlap those of the NiFe2 O4 phase. It must also be noted that the stoichiometry of the nickel ferrite phase is approximate because variation in its composition induces only very small shifts in the XRD peak positions [19] that are difficult to assess in the present case. Regarding the Cu-free alloy, the surface scale is mainly composed of Fe2 O3 . Fig. 4 shows the cross-sectional SEM images and corresponding EDX maps of O, Ni, Fe and Cu elements for the oxidized (after TGA) Cu-65, Cu-55, Cu-45 and Cu-0 samples. This is complemented by EDX analyses performed in different zones of the cross-sectioned samples as shown in Fig. 5 A–D. As support to the TGA results, the thickness of the surface oxide scale on the Cu–Ni–Fe alloys decreases with their Cu content (∼35, 10 and 7 ␮m for Cu-65, Cu-55 and Cu-45, respectively). The oxide scales are rather dense, except for the Cu-65 sample in which a few macropores are apparent. The oxide scale formed on the three Cu–Ni–Fe samples presents a bilayer structure. The outer layer is exclusively constituted of Cu oxides. Its thickness decreases with

the alloy Cu content (∼20, 5 and 3 ␮m for Cu-65, Cu-55 and Cu-45, respectively). On the basis of the Cu/O atomic ratio determined by EDX analyses, the outer layer is formed of CuO in all three cases. The inner oxide layer of Cu-55 and Cu-45 sample is almost exclusively composed of Fe and Ni. These two elements are homogeneously distributed in the inner oxide layer with a Fe/Ni atomic ratio close to 2 as revealed from EDX analysis (see for instance zone 3 in Fig. 5B and zone 5 in Fig. 5C), which tends to confirm the formation of NiFe2 O4 in this layer. For the Cu-65 sample, the inner oxide layer is much thicker (∼15 ␮m compared to ≤5 ␮m for the Cu-55 and Cu-45 samples). Moreover, it contains a significant amount of Cu oxides with a decreasing concentration gradient towards the oxide/alloy interface (see zones 3–5 in Fig. 5A). A Cuenriched region (∼10 ␮m thick) is clearly discernable in the Cu-45 alloy just below the oxide scale as highlighted in Fig. 4 (see Cu map) and in Fig. 5C (see zone 6). This Cu-enriched region is less marked for the Cu-55 alloy and absent for the Cu-65 alloy. One can also note the presence of Fe oxide inclusions in the bulk alloy (see for instance zone 9 in Fig. 5A), which are more numerous and deeper in the Cu-65 alloy than in the Cu-55 and Cu-45 alloys. The structure of the oxide scale observed on the Cu–Ni–Fe samples in Figs. 4 and 5A–C can be explained as follows. Cu, Ni and Fe elements are oxidized simultaneously at the initial stage of oxidation. A gradient of oxygen activity is established gradually in the oxide scale, which drives the elements to diffuse in the oxide scale. Finally, CuO with the highest equilibrium oxygen pressure forms at the top of composite oxide scale to keep the balance on thermodynamics. According to the Cu–O stability diagram [20], the equilibrium O2 pressure for CuO/Cu2 O is around 10−4 atm at 700 ◦ C. Accordingly, CuO is formed next to the sample surface whereas Cu2 O is stable inside the oxide scale, as shown for the Cu-65 sample. For the Cu-55 and Cu-45 samples, the outer layer is thinner

108

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

Fig. 4. Cross-sectional SEM images and corresponding EDX maps of O, Ni, Fe and Cu elements for the oxidized (after TGA) Cu-65, Cu-55, Cu-45 and Cu-0 samples.

due to a lower outward flux of Cu atoms that are almost fully converted into a single CuO phase (i.e. the rate of complete oxidation of the copper is faster than the rate of copper diffusion for Cu-55 and Cu-45 samples). As a consequence of the outward Cu diffusion, a Cu-depleted region is formed underneath. Oxygen atoms can diffuse into that Cu-depleted region and internal oxidation of Fe and Ni can occur to form NiO and Fe2 O3 , which react together to form a continuous inner NiFe2 O4 layer (G◦ f (1000 K) = −727 kJ mol−1 )[21], as shown for the Cu-55 and Cu-45 samples. In the case of the Cu-65 alloy, the NiO and Fe2 O3 particles are dispersed in a CuOx matrix and less able to interact together and form a continuous and dense NiFe2 O4 layer in the inner region. Since NiFe2 O4 acts as a diffusion barrier layer at the oxide-alloy interface, the oxidation kinetics of Cu-55 and Cu-45 alloys are slowed down as demonstrated by the TGA curves (Fig. 2). Its efficiency in blocking the outward Cu diffusion is also highlighted by the formation of a Cu-enriched region underneath the NiFe2 O4 layer for the Cu-55 and Cu-45 samples. As mentioned earlier, a continuous NiFe2 O4 layer is not formed in the Cu-65 sample. Accordingly, oxygen diffusion in the bulk alloy can occur leading to the oxidation of the less noble element (Fe) of the Cu–Ni–Fe alloy, as confirmed by the presence of Fe oxide inclusions in the bulk alloy. In the case of the Cu-55 and Cu-45 samples, it is surmised that oxidation diffusion is slowed down by the presence of the NiFe2 O4 layer since these Fe oxide inclusions in the bulk alloy are much less numerous. However, as seen in Figs. 5, defects such as cracks and pores are formed in the oxide scale, which must lead oxygen to transport to the oxide-alloy interface. Therefore, for longer oxidation times, oxidation of Cu underneath the NiFe2 O4 layer may proceed, and a new sequence in the layered CuOx /NiFe2 O4 structure could be initiated as observed by Haugsrud et al. on as-cast (biphased) 55Cu–30Ni–15Fe alloy oxidized at 750 ◦ C in 1 atm O2 for 760 h

[22]. It will be relevant to perform longer oxidation treatment (»20 h) with the present Cu–Ni–Fe alloys to confirm this issue and to evaluate the impact of their monophased structure on this process. In comparison, as shown in Figs. 4 and 5D, the oxide scale formed on the Cu-0 sample is mainly constituted of Fe oxides near the surface (as Fe2 O3 on the basis of XRD analysis, Fig. 3). The thickness of the oxide scale is not uniform (10–15 ␮m) with a dendritic morphology at the oxide/alloy interface. The preferential outward Fe oxidation results in a Ni-enriched region in the alloy next to the oxide/alloy interface. Ni is also detected in the oxide scale but it is not homogeneously distributed in the oxide scale as evidenced by the presence of large Ni(Fe)O inclusions in the inner part of the oxide layer (see zone 4 in Fig. 5D). EDX data (see zone 2 in Fig. 5D) suggest the formation of NiFe2 O4 between the outer Fe2 O3 layer and the inner Ni(Fe)O-rich scale. However, the absence of peaks relative to this phase in the XRD pattern (Fig. 3) tends to indicate that NiFe2 O4 is not extensively formed during the alloy oxidation in comparison to the previous Cu–Ni–Fe alloys. This means that Cu in appropriate proportions in Cu–Ni–Fe alloys favors the formation of the NiFe2 O4 phase. A possible explanation is that the CuOx layer formed at the surface of the Cu–Ni–Fe alloys during the first stage of the oxidation acts as a “protective layer” which may give time for the internal oxidation of Fe and Ni to occur and also favor the subsequent formation of a coherent inner NiFe2 O4 scale. At the extreme, without this surface layer, the rapid surface segregation of Fe2 O3 during the oxidation of the present Ni–Fe alloy results in the formation of an inner Fe-depleted region with a Ni/Fe ratio less favourable for the formation of NiFe2 O4 . This is in accordance with the study of Chapman et al. on the oxidation of Ni–Fe alloys which showed that the formation of NiFe2 O4 strongly depends on the concentration of Ni at the oxide/alloy interface [23].

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

109

3.3. Electrochemical behavior for Al electrolysis Fig. 6 compares the evolution of the cell voltage for 20 h of electrolysis at an anodic current density of 0.5 A cm−2 in aluminasaturated KF–AlF3 electrolyte at 700 ◦ C with the Cu-65, Cu-55, Cu-45 and Cu-0 anodes. For the Cu-65 anode, a progressive increase of the cell voltage is observed during the first 13 h of electrolysis. This is attributed to the growth of the oxide scale on the anode decreasing the anode electrical conductivity, which results in an increase of the ohmic drop component of the cell voltage [12]. After this period, the cell voltage nearly stabilizes aroung 4.7 V, suggesting that a steady state in the oxide formation/dissolution processes is reached. A similar behaviour is observed for the Cu-55 anode. In contrast, for the Cu-45 and Cu-O anodes, an abrupt and large increase of the cell voltage is observed after 5–10 h of electrolysis, which could be attributed to the formation of an insulating scale onto these anodes. The brief potential drops observed during the last hours of electrolysis is probably due to sudden spalling of the surface oxide layer under the action of the O2 bubbles, reflecting its poor mechanical stability (to be confirmed).

Fig. 5. (Continued ).

Fig. 5. EDX analyses in different zones of the oxidized Cu-65 (A), Cu-55 (B), Cu-45 (C) and Cu-0 (D) samples.

Fig. 7 shows cross-sectional photographs of the 20 helectrolyzed anodes embedded-polished in epoxy resin. The initial size of the anodes is represented by the red dashed lines. For the Cu–Ni–Fe anodes, an increase of their thickness of ∼15% is observed whereas a decrease of ∼20% is seen for the Ni–Fe anode. This highlights the much higher dissolution rate of the copper-free alloy. This is also confirmed by the lower purity of the produced Al with the Ni–Fe anode (92% compared to >99% with the Cu–Ni–Fe anodes). It is thus evident that Cu plays a key role in the corrosion resistance of the Cu–Ni–Fe anodes. However, it must be remembered that a too high Cu content in the alloy induces a major increase of the Cu contamination in the produced Al [10]. Cross-sectional BSE images and corresponding elemental EDX maps and profiles (F, Al, K, O, Ni, Fe, Cu elements) of the 20 helectrolyzed Cu-65, Cu-55, Cu-45 and Cu-0 anodes are presented in Fig. 8. For the three Cu-containing electrodes, the surface oxide layer has a thickness of about 500–600 ␮m, which is much thicker than on the oxidized samples (7–35 ␮m). A partial delamination is observed at the oxide/alloy interface, probably because of the thermal shock during its extraction from the electrochemical reactor.

110

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

Fig. 5. (Continued ). Fig. 5. (Continued ).

It may also be explained by lacunar coalescence creating porosity due to Cu diffusion towards the bath. On the basis of XRD analyses (not shown), it is constituted in all cases of CuOx , NiFe2 O4 and KF–AlF3 electrolyte. The presence of NiO and FeOx is also possible. No well-defined layered structure appears in the oxide scale in contrast to the oxidized samples. A succession of non-continuous Ni–Fe-rich and Cu-rich oxide layers is however discernable for the Cu-45 anode. The main difference between the three Cu–Ni–Fe anodes is that electrolyte penetration in the oxide layer is strongly accentuated with the decreasing Cu content in the alloy, as clearly evidenced by the F, Al and K EDX mapping images in Fig. 8. The very high electrolyte penetration for the Cu-45 anode can explain its large cell voltage variation observed in Fig. 6. The accentuation of the electrolyte penetration with the decrease of the Cu content in the alloy is rather unexpected. Indeed, a lower Cu content in the alloy should induce a lower outward Cu diffusion and thus a less porous Cu-depleted region underneath for electrolyte penetration. Moreover, on the basis of the previous oxidation experiments, the formation of a denser and thus more protective inner NiFe2 O4 layer is expected for the Cu-55 and Cu-

45 anodes. It is hypothesized that the presence of CuO in the outer layer is critical for the formation on an inner layer of NiFe2 O4 . When the electrode is in contact with the KF–AlF3 electrolyte at 700 ◦ C, the outer CuO layer is readily dissolved. Therefore, the reservoir of Cu must be large enough for the outward flux of Cu from the alloy to counterbalance its dissolution. This is the case for the Cu-65 sample and the outer CuO layer acts as a sacrificial layer that allows enough time for the formation of the NiFe2 O4 inert layer to occur. This is not the case for the less Cu-rich anodes (Cu-55 and Cu-45 samples). This tends to be confirmed by the EDX profiles (Fig. 8), showing the absence of Cu in the outermost layer (first ∼50 ␮m) for the Cu-55 and Cu-45 anodes in contrast to the Cu-65 anode. For the Cu-0 electrode, the surface scale is much thinner (∼100 ␮m) and contains a large amount of electrolyte. NiO and FeOx are also detected in the surface scale. The presence of NiFe2 O4 cannot be excluded on the basis of the EDX and XRD (not shown) analyses. However, considering the very high dissolution rate of this anode, it can be assumed that the formation of protective (low soluble) NiFe2 O4 phase is limited with the Cu-free anode as also shown for the oxidized Ni–Fe sample. This confirms that Cu in Cu–Ni–Fe alloys plays a major role in the formation of the NiFe2 O4

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

7

7

Cu-65

Cu-55 6 Cell voltage (V)

Cell voltage (V)

6

5

4

5

4

3 0

5

10

15

3

20

0

5

10

15

20

15

20

Electrolysis time (h)

Electrolysis time (h)

7

7

Cu-45

Cu-0

6

6 Cell voltage (V)

Cell voltage (V)

111

5

4

5

4

3

3 0

5

10

Electrolysis time (h)

15

20

0

5

10

Electrolysis time (h)

Fig. 6. Cell voltage vs. electrolysis time at Ianode = 0.5 A cm2 in alumina-saturated KF–AlF3 electrolyte at 700 ◦ C for the Cu-65, Cu-55, Cu-45 and Cu-0 anodes.

Fig. 7. Cross-sectional photographs of the 20 h-electrolyzed anodes embedded-polished in epoxy resin. The initial size of the anodes is represented by the red dashed lines.

112

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

Fig. 8. Cross-sectional BSE images and corresponding elemental EDX maps and profiles (F, Al, K, O, Ni, Fe and Cu elements) for the 20 h-electrolyzed Cu-65, Cu-55, Cu-45 and Cu-0 anodes.

E. Gavrilova et al. / Corrosion Science 101 (2015) 105–113

phase by preventing the rapid Fe oxide surface segregation as previously discussed. 4. Conclusion This work has provided strong support for a mechanism of both reaction and protection of Fe, Cu, Ni containing inert anodes for aluminum production. Due to its relatively high mobility, copper appears to move preferentially to the outer surface of the anode. The copper at the surface reacts with oxygen, and minimizes oxygen transport into the bulk alloy. The region that is depleted of copper gives the right conditions for the formation of a nickel spinel phase (NiFe2 O4 ) which retards further movement of copper to the surface and also resists oxygen diffusion. The NiFe2 O4 phase also decreases electrolyte penetration in the oxide scale and prevents the rapid anode dissolution. The Cu concentration in the alloy is critical for the formation of the spinel, since too much copper prevents the spinel from forming as a coherent inner layer and too little does not give the required time to allow the spinel to form without excess oxygen present. Acknowledgements The authors thank the Natural Sciences and Engineering Research Council of Canada (NSERC) and Kingston Process Metallurgy (KPM) Inc. for supporting this work. References [1] ASME Technical Working Group, Inert anode technologies report, U.S. Department of Energy, Office of Industrial Technologies, Washington, 1999. [2] I. Galasiu, R. Galasiu, J. Thonstad, Inert Anodes for Aluminium Electrolysis, Aluminium-Verlag, Dusseldorf, 2007. [3] R.P. Pawlek, Inert anodes: an update, Light Met. (2014) 1309–1313. [4] J.N. Yang, B.R. Hryn, A. Davis, G.K. Roy, New opportunities for aluminum electrolysis with metal anodes in a low temperature electrolyte system, Light Met. (2004) 321–326. [5] D. Yang, C. Graczyk, Alumina solubility in KF-AlF 3-based low-temperature electrolyte system, Light Met. (2007) 537–541.

113

[6] L. Cassayre, P. Palau, P. Chamelot, L. Massot, Properties of low-temperature melting electrolytes for the aluminum electrolysis process: a review, J. Chem. Eng. Data 55 (2010) 4549–4560. [7] T.R. Beck, R.J. Brooks, Non-consumable anode and lining for aluminium electrolytic reduction cell, US patent 5,284,562 (1994). [8] T.R. Beck, A non-consumable metal anode for production of aluminum with low-temperature fluoride melts, Light Met. (1995) 355–360. [9] T.R. Beck, C.M. MacRae, N.C. Wilson, Metal anode performance in low-temperature electrolytes for aluminum production, Metall. Mat. Trans. B 42 (2011) 807–813. [10] S. Helle, M. Pedron, B. Assouli, B. Davis, D. Guay, L. Roué, Structure and high-temperature oxidation behavior of Cu–Ni–Fe based alloys prepared by high-energy ball milling for application as inert anodes in aluminum electrolysis, Corros. Sci. 52 (2010) 3348–3355. [11] G. Goupil, S. Jucken, D. Poirier, J.G. Legoux, E. Irissou, B. Davis, D. Guay, L. Roué, Cold-sprayed Cu–Ni–Fe anode for Al electrolysis, Corros. Sci. 90 (2015) 259–265. [12] S. Helle, M. Tresse, B. Davis, D. Guay, L. Roué, Mechanically alloyed Cu–Ni–Fe–O based materials as oxygen-evolving anodes for aluminum electrolysis, J. Electrochem. Soc. (2012) E62–E68. [13] G. Goupil, S. Helle, B. Davis, D. Guay, L. Roué, Anodic behavior of mechanically alloyed Cu–Ni–Fe and Cu–Ni–Fe-O electrodes for aluminum electrolysis in low-temperature KF–AlF3 electrolyte, Electrochim. Acta 112 (2013) 176–182. [14] S. Helle, B. Brodu, B. Davis, D. Guay, L. Roué, Influence of the iron content in Cu–Ni based inert anodes on their corrosion resistance for aluminum electrolysis, Corros. Sci. 53 (2011) 3248–3253. [15] S. Helle, B. Davis, D. Guay, L. Roué, Electrolytic production of aluminum using mechanically alloyed Cu–Al–Ni–Fe based materials as inert anodes, J. Electrochem. Soc. 157 (2010) E173–E179. [16] H.X. Li, X.J. Hao, G. Zhao, S.M. Hao, Characteristics of the continuous coarsening and discontinuous coarsening of spinodally decomposed Cu–Ni–Fe alloy, J. Mater. Sci. 36 (2001) 779–784. [17] M. Baricco, E. Bosco, G. Acconciaioco, P. Rizzi, M. Coisson, Rapid solidification of Cu–Ni–Fe alloys, Mater. Sci. Eng. A 375–377 (2004) 1019–1023. [18] I. Gallino, M.E. Kassner, R. Bush, Oxidation and corrosion of highly alloyed Cu–Ni–Fe as inert anode material for aluminium electrowinning in as-cast and homogenized conditions, Corros. Sci. 63 (2012) 293–303. [19] S. Corso, P. Tailhades, I. Pasquet, A. Rousset, V. Laurent, A. Gabriel, C. Condolf, Preparation conditions of pure and stochiometric Nix Fe3-x O4 bulk ceramics, Solid State Sci. 6 (2004) 791–798. [20] ASM handbook, Volume 3: Alloy Phase Diagrams, ASM Publishing, 1992, pp 2•174. [21] JANAF Thermochemical Tables, 3rd ed., Vol. 14, U.S. National Bureau of Standards, Washington, DC (1985). [22] R. Haugsrud, T. Norby, P. Kofstad, High-temperature oxidation of Cu–30 wt%Ni–15 wt% Fe, Corros. Sci. 43 (2001) 283–299. [23] V. Chapman, B.J. Welch, M. Skyllas-Kazacos, Anodic behaviour of oxidised Ni–Fe alloys in cryolithe–alumina melts, Electrochim. Acta 56 (2011) 1227–1238.