Electrochemical behavior of a quasi-crystalline AlPdMn alloy in a chloride-containing solution

Materials Science and Engineering, A 181/14182 (1994) 1141 - 1144

1141

Electrochemical behavior of a quasi-crystalline A1-Pd-Mn alloy in a chloride-containing solution K. Asami, A.-P. Tsai and K. Hashimoto Institute for Materials Research, Tohoku University, Sendai 980 (Japan)

Abstract The electrochemical properties of a stable quasi-crystalline A I - P d - M n alloy were studied by means of the potentiodynamic polarization method in a deaerated NaC1 solution at 30 °C. The surface morphology and composition after polarization in the solution were observed by scanning electron microscopy combined with energy-dispersive X-ray analysis. The results were compared with those of a crystalline A1-Pd-Mn alloy and the constituent elements A1, Mn and Pd. The quasi-crystalline A1-Pd-Mn alloy cracked into small pieces during corrosion in the above solution. The cracking phenomenon was attributed to the hydrogen embrittlement caused by the combination of hydrogen evolved by the corrosion of the alloy and the internal tensile stress of the alloy. After the corrosion, both alloys were enriched in Pd on their surfaces.

1. Introduction Since the existence of quasi-crystals was made clear by Shechtman et al. [1] and Levine and Steinhardt [2] in 1984, much effort has been devoted to the investigation of quasi-crystalline alloys. However, most of the research work on quasi-crystalline alloys has been directed at structural analysis, and surveying of new quasi-crystalline alloy systems and their physical properties. Following the discovery of a stable quasicrystalline alloy of A156Li33Cull by Sainfort and Dubost [3] in 1986, some of the present authors have discovered a number of stable quasi-crystalline alloy systems [4, 5]. To our knowledge, there are few reports on the chemical properties of quasi-crystalline alloys. Recently, Massiani et al. [6] studied the chemical properties of several Al-based alloys containing a quasi-crystalline phase, as a major constituent or approximant phase, in NaOH + H 2 S O 4 solutions. Since quasi-crystals have very high hardness values, they can be applied as composite materials, in combination with other structural materials. For practical use of such materials, however, it is necessary to know their chemical properties in environments in which they are to be exposed. In this paper, some electrochemical properties of a quasi-crystalline AI-Pd-Mn alloy were investigated and compared with those of a crystalline A1-Pd-Mn alloy and the constituent elements in a solution containing chloride ions (which have the 0921-5093/94/$7.00 SSD1 0921-5093(93)05648-9

largest detrimental corrosive effect on ordinary metallic materials).

2. Experimental details Two AI-Pd-Mn alloys, A172Pd20Mn 8 and A160Pd25Mnl5 , were prepared by the Ar arc melting method. To homogenize the alloys, the alloy ingots were annealed at 1073 K (800 °C) for 12 h in vacuo. Examination of the crystallinity and the phase identification of the alloys thus prepared were carried out by X-ray diffraction (XRD). The electrochemical properties of the alloys were studied in a deaerated borate-boric acid buffer solution (pH 8.4) containing 0.5 M NaCI at 303 K (30°C), by measuring the potentiodynamic polarization curves at a potential scanning rate of 50 mV min- l, as well as the open circuit potential (OCP) measurement. When the initial OCP had nearly reached the steady state, the polarization curve measurements were started. The change in the OCPs with time was also traced immediately after the polarization measurements. The potentials were all referred to the saturated calomel electrode (SCE). For comparison, the polarization curves and OCPs of pure Pd, Mn and AI were also measured in the above solution under the same conditions. © 1994 - Elsevier Sequoia. All rights reserved

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Electrochemical behavior of Al-Pd-Mn alloy

The surface morphology and composition were observed using a scanning electron microscope (Hitachi $800) equipped for energy-dispersive X-ray (EDX) analysis. The EDX results were analyzed by a standardless method.

3. Results and discussion The powder XRD patterns of the fully annealed A172Pd20Mn8 and A160Pd25Mn,5 alloys are shown in Fig. 1. (The subscript number attached to each atomic symbol indicates the atomic percentage of the element within the alloy.) All the diffraction peaks of the Al7~Pd20Mn 8 alloy can be indexed as an icosahedral quasi-crystalline phase, although not all the indexes are indicated in the figure. It can be said that the A172Pd20Mn8 alloy consists of a single icosahedral quasi-crystalline phase. It can also be understood that the quasi-crystallinity of the A172Pd~oMn 8 alloy thus prepared is very high, since the peak separation caused by the Cu K a I and K a : doublet is clearly observed on peaks at 2 0 angles higher than 50 °. The main diffraction peaks of the A160Pd25Mn15 alloy can be indexed by a cubic phase. There remain several peaks corresponding to unknown metallic compound(s). It follows that the Al60Pd2~Mnl~ alloy can be a mixture of a cubic phase and unknown metallic compound(s).

The OCPs of both the alloys are shown in Figs. 2 and 3. For comparison, those for pure Pd, Mn and A1 are also illustrated in Fig. 4. The initial OCPs of both the quasi-crystalline A172Pd20Mn8 (Q alloy) and the crystalline A160Pd25Mn~5 (C alloy) are much higher than those of Mn and AI, but lower than that of Pd. The OCP of AI increases at first but then decreases after about 100 min, reaching about 1.2 V (SCE) after several hundred minutes. The initial OCPs of both the Q and C alloys oscillate randomly around - 0 . 4 5 V (SCE), although the OCP of the Q alloy tends to be slightly more positive. The OCPs of the Q and C alloys after the anodic and cathodic polarization measurements are also inserted in Figs. 2 and 3. Both alloys show almost the same OCP after the polarize-

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Electrochemical behavior of Al-Pd-Mn alloy ......

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tion, and converge to about - 0 . 6 5 V (SCE) without potential oscillation. The OCPs of both the Q and C alloys after the polarization are about 0.2-0.3 V less positive than the initial values. The difference between the initial OCP and that after the polarization could arise from the change in the alloy surface composition caused by the preferential dissolution of AI, as will be mentioned later. The existence of the air-formed films on the alloys also could play a role, although they do not have high corrosion resistance, because the initial OCPs of these alloys show random oscillation with time, indicating the breakdown and reformation of the surface film in the solution. The anodic and cathodic polarization curves of the Q and C alloys are shown in Fig. 5. The corresponding polarization curves for the constituent elements are

also shown in Fig. 6 for comparison. The anodic polarization curve of the Q alloy shows spontaneous passivation, though it ends up with a pitting-like sudden current increase in the anodic current density on further anodic polarization, while the anodic current density of the C alloy increases abruptly from the O C E The cathodic activity of the Q alloy is as high as that of pure Pd near the OCP region, while that of the C alloy is low. It can be seen that the cathodic polarization curve of the C alloy coincides with that of pure Mn, probably because of the high Mn content of the C alloy. The scanning electron microscopy (SEM) images of the surface morphology of the Q and C alloys after the anodic polarization are shown in Fig. 7. Corresponding to the XRD results, the SEM image of the as-polished C alloy had two types of grain: one type had a darker contrast than that of the other type. The dark grains seem to have been preferentially dissolved by the anodic polarization. The SEM image of the Q alloy surface in Fig. 7 shows that the surface is cracked into small pieces by the anodic polarization, and that the crack mouths are wide open. There is no sign of morphology corresponding to ordinary pitting corrosion. It was observed that bubbles evolved vigorously from the surface during anodic polarization when the anodic current was high. Since the potential of the gas evolution observed was far lower than the 02 and C1z evolution potentials, this phenomenon is analogous to the amorphous Ni-Zr alloys in a hydrochloric solution, where H 2 evolves by the direct reaction of the surface with the solution, as a result of the breakdown of the surface film by the anodic polarization [7]. Similarly to the amorphous Ni-Zr alloys, hydrogen embrittlement may have occurred under the anodic polarization conditions. In the case of the amorphous

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K. Asami et al.

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Electrochemical behavior of Al-Pd-Mn alloy

TABLE 1. Results of a semi-quantitative EDX analysis of the average composition (at.%) of the quasi-crystalline and crystalline alloys before and after anodic polarization

As-polished After anodic polarization

Quasi-crystalline A172Pd20Mn8

Crystalline A160Pd25Mn15

AI

Pd

Mn

A1

Pd

Mn

69.7 30.5

21.0 61.0

9.2 8.4

56.4 40.5

25.3 0.5

18.3 13.7

depleted after the anodic polarization. This might well be caused by the preferential dissolution of AI. It is found by XRD that the C alloy consists of two or more phases. The SEM observation reveals that there are two types of grain, i.e. grains with dark and bright contrast: one grain type is relatively rich in both AI and Mn, and the other is rich in Pd. Therefore, it can be said that the former phase must be preferentially corroded in the solution examined. 4. Conclusions

Fig. 7. Scanning electron micrographs of (a) quasi-crystalline A172Pd20Mn8 and (b) crystalline A160Pd25Mn15alloy after anodic polarization in the deaerated borate-boric buffer solution (pH 8.4) containing 0.5 M NaC1 at 303 K (30 °C).

alloys, the cracks developed in the radial direction from the initiation spot, because of the uniformity of the inner stress. For the Q alloy, the cracks run mainly in the direction of the mechanical polishing. However, there are also many cracks running across the cracks parallel to the polishing scratch marks. There should certainly exist some intrinsic tensile stress in the Q alloy. The C alloy was similarly polished under the same conditions, yet it did not have the susceptibility to the hydrogen embrittlement. It seems, therefore, that the tensile stress is intrinsic to the Q alloy and that the cracking p h e n o m e n o n is attributable to the hydrogen embrittlement caused by the combination of H 2 evolved by the corrosion of the alloy and the internal tensile stress of the Q alloy. T h e E D X results are given in Table 1. It can be clearly seen that, on average, Pd is enriched and AI is

Crystalline A160Pd25Mn~5 alloy consists of two phases: one is relatively rich in Pd and the other one is rich in both Mn and A1. The second phase is preferentially corroded in a deaerated borate-boric buffer solution (pH 8.4) containing 0.5 M NaC1 at 303 K (30 °C). The single-phase quasi-crystalline A172Pd20Mn 8 alloy, which has an internal tensile stress, cracked into small pieces during corrosion in the above solution. The cracking phenomenon is attributable to the hydrogen embrittlement caused by the combination of H 2 evolved by the corrosion of the specimen and the internal tensile stress of the quasi-crystalline alloy. After the corrosion, both alloys are enriched in Pd on their surfaces. References 1 D. Shechtman, I. Blech, D. Gratias and J. W. Cahn, Phys. Rev. Lett., 54 (1984) 1951. 2 D. Levine and P. J. Steinhardt, Phys. Rev. Lett., 54 (1984) 2477. 3 P. Sainfort and B. Dubost, J. Phys. (Paris), Colloq. C3, (1986) C3-321. 4 A. P. Tsai, A. Inoue and T. Masumoto, Jpn. Appl. Phys., 26 (1987) L1505; 27 (1988) L1587. 5 A.P. Tsai, A. Inoue and T. Masumoto, Mater Trans. Jpn. Inst. Met., 30 (1989) 300,463. 6 Y. Massiani, A. Ait Yaazza, J. P. Crousier and J. M. Dubois, J. Non-Cryst. Solids, 159 (1993) 92. 7 K. Asami, H. Habazaki, K. Kawashima and K. Hashimoto, Corros. Sci., 34 (1993) 445.