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The corrosion behaviour of sputter-deposited amorphous NiTi alloys in 1 M HCl
Materials Science and Engineering, A 181/,4182 (1994) 1128-1132
1128
The corrosion behaviour of sputter-deposited amorphous Ni-Ti alloys in 1MHC1 Eiji Akiyama, Hiroki Habazaki, Asahi Kawashima, Katsuhiko Asami and Koji Hashimoto Institute for Materials Research, Tohoku University, Sendai 980 (Japan) Abstract Amorphous Ni-Ti alloys were prepared by d.c. magnetron sputtering. The corrosion behaviour of amorphous Ni-Ti alloys in 1 M HCI was investigated by electrochemical measurements. Particular attention was paid to the compositions of the surface film and underlying alloy surface just below the surface film. The amorphous Ni-Ti alloys were spontaneously passive. An X-ray photoelectron spectroscopy analysis revealed that the thickness of the surface film was increased linearly with an increase in the polarization potential. The growth rates for amorphous Ni-32Ti and Ni-62Ti alloys were the same as each other and slightly lower than the growth rate of sputter-deposited titanium. The passive film formed on the sputter-deposited amorphous Ni-Ti alloys was enriched in titanium cation. In the underlying alloy surface just below the surface film after polarization for 1 h, the titanium content increased linearly with the increase in the potential. At relatively low potentials, nickel was concentrated in the underlying alloy surface, while at relatively high potentials the content of titanium was higher than that in the bulk alloy.
1. Introduction
It has been found that the corrosion resistance of melt-spun amorphous Ni-Ti alloys is higher than that of the crystalline alloys [1], and that melt-spun amorphous Ni-Ta and Ni-Nb alloys have a high corrosion resistance even in hot concentrated HNO 3 [2], HC1 [2, 3] and H 3 P O 4 [4, 5].
However, the practical applications of the melt-spun amorphous alloys are limited by their shape. By contrast, sputer deposition is suitable for the preparation of the amorphous alloys with a wide surface area on conventional crystalline bulk metals. It is expected that the amorphous Ni-valve metal alloys prepared by sputtering also have high corrosion resistance. The present work aims to prepare corrosionresistant Ni-Ti alloys by a d.c. magnetron sputter deposition method and to clarify their corrosion and anodic polarization behaviour in 1 M HC1 solution. For a better understanding of the corrosion resistance of the alloys, their surface structures were analysed. Particular attention was paid to the compositions of the surface film and underlying alloy surface just below the surface film.
2. Experimental method
D.c. magnetron sputtering was carried out for preparation of Ni-Ti alloys. The target was composed of a 99.95% pure titanium disc of 100 mm diameter 0921-5093/94/$7.00 SSD10921-5093( 9 3 )05640-B
and 6 mm thickness, on the sputter erosion region of which 99.95% pure nickel discs of 20 mm diameter and 10 mm diameter were placed. The composition of the sputter deposits was changed by changing the numbers of nickel discs on the titanium disc. Glass plates were used as substrates. The sputtering apparatus and conditions used were the same as those described elsewhere [5, 6]. For the purpose of homogenization of the sputter deposits, the watercooled substrates were revolved around a central axis of the sputtering chamber, in addition to revolution of the substrates themselves around their own axes. After the target and substrates were installed in the sputtering machine the vacuum chamber was evacuated to about 1 x 10-6 Torr. After presputtering of the target for 5-10 min, sputtering was carried out at (4-8) x 10-4 Torr of argon gas. The composition of the alloys prepared was determined by inductively coupled plasma emission spectrometry (ICP) after the alloys were dissolved into 6 M HC1 solution. Their structure was identified by X-ray diffraction with Cu K a radiation. Potentiostatic polarization curves were measured in a deaerated 1 M HCI solution at 30 °C. In order to obtain reproducible results removal of the oxide film formed on as-sputtered alloys by air exposure was necessary, and hence, before polarization or immersion, the specimens were polished mechanically with silicon carbide paper up to grit 1500 in cyclohexane and dried in air. A platinum electrode and a saturated calomel elecrode (SCE) were used as counterelectrode © 1994 - Elsevier Sequoia. All rights reserved
E. Akiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni-Ti
and reference electrode respectively. Potentials hereafter are referred to the SCE. After potentiostatic polarization, X-ray photoelectron spectra were measured by means of a Shimadzu-850 photoelectron spectrometer with Mg K a radiation. For quantitative analysis of the alloy surface, the photoionization cross-section of the Ni 2p3/2 electrons relative to the O ls electrons including the satellite of the Ni 2p3/2 spectrum was determined by using the standard substances Ni(OH)2 and Ni203 as 2.315. The photoionization cross-section of the Ti 2p3/2 electrons relative to O ls electrons used was 1.277 [7]. The thickness and composition of the surface film and the composition of the underlying alloy surface were quantitatively determined by a previously proposed method using integrated intensities of photoelectrons [8]. A solution analysis was performed by means of ICE The concentrations of nickel and titanium dissolved were measured at 231.604 nm and 334.941 nm wavelengths respectively.
range. The broad peak at about 23 ° arises from the glass substrate. When the alloy titanium content increases, the continuous shift of the first halo position was observed from that close to the 111 reflection of the f.c.c, nickel phase to that close to the 011 reflection of the h.c.p, titanium phase. This indicates that the alloy consists of a single amorphous phase. Figure 2 shows the potentiostatic polarization curves for the amorphous Ni-32Ti and Ni-62Ti alloys. The amorphous Ni-Ti alloys are spontaneously passive and their anodic current densities are almost constant in the passive region. An abrupt increase of the current density is observed at about 1.5 V. This current increase is attributable to oxygen evolution. The current density of the Ni-32Ti alloy is higher than that of the Ni-62Ti alloy. Accordingly the corrosion resistance of the Ni-62Ti alloy is higher than that of the Ni-32Ti alloy. Figure 3 shows the change in the thickness of the surface film formed on the amorphous Ni-32Ti and Ni-62Ti alloys and sputter-deposited titanium as a
10 3
3. Experimental results
l
101
I
//
Ti sputtered
=- 10 0 e,-
[
l
o;c
'E
a
I
deaerated 1 M HCI
¢q
Figure 1 shows X-ray diffraction patterns of sputterdeposited Ni-Ti alloys. Alloy compositions hereafter are all denoted in atomic percentages. When the titanium content is very low, such as in Ni-16Ti alloy, reflections close to the 111,200 and 220 reflections of the f.c.c, nickel phase are observed. The alloys containing 32 at.% or more of Ti show the halo pattern typical for the amorphous structure over a wide composition
1129
//
....\
10 1
//
Ni-32Ti
:~
f
= 10-2
o 10 -3
I
-1.0
I
-0.5
0.0
0.5
1.0
1.5
2.0
P o t e n t i a l / V v s SCE
I
I
I
I
I
Fig. 2. Potentiostatic polarization curves of sputter-deposited amorphous Ni-32Ti and N i - 6 2 T i alloys in deaerated 1 M HC1 at 30 °C.
I
70
I
I
I
deaerated 1M HCI o
8 C
50
._o .c= I-- 40 IE 30
as-polished
~
-1.0
Ni-7,1Ti
I
Ni-62Ti
2Ti
qr 20
I
___~ -~//'~~
" 13
30 C, l h
,~ 6O
Ti sputtered
Ni-62Ti : 1 6 ~ v '
J
I
I
I
-0.5
0.0
0.5
1.0
I
1.5
2.0
Potential / V vs SCE
20
I
I
I
i
I
I
30
40
50
60
70
80
90
2 0 / degree Fig. 1. X-ray diffraction patterns of sputter-deposited Ni-Ti alloys.
Fig. 3. Thickness of surface films formed on the amorphous Ni-32Ti and Ni-62Ti alloys and sputter-deposited titanium in 1 M HC1 at 30 °C. The thickness of the air-formed film after mechanical polishing in cyclohexane is also shown for comparison.
1130
E. Akiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni-Ti
function of polarization potential. This figure clearly reveals that anodic polarization leads to film thickening with the polarization potential. The apparent growth rates calculated from the slope of the straight lines for the amorphous Ni-32Ti and Ni-62Ti alloys are both approximately 16 A V- 1. The growth rate for sputterdeposited titanium is 21 A W-1, which is almost the same as the anodizing ratio reported for titanium [9]. The growth rate for the amorphous Ni-Ti alloys is smaller than that for sputtered titanium, perhaps because the surface film contains nickel ions. Figures 4 and 5 show the cationic fraction of titanium in the passive film and the atomic fraction of titanium in the underlying alloy surface just below the passive film for the amorphous Ni-32Ti and Ni-62Ti alloys after anodic polarization for 1 h as a function of Ti
E 1.0 i,
= E 0.8 _= e= ¢0
~- 0.6
"6 tO
0.4
O [] z~ •
o
.,= U.
_o 0.2 e. o
o 0.0
deaerated 1M HCI 30°C I
I
I
Ni-32Ti, lh Ni-32Ti, 2h Ni-32Ti, 4h Ni-62Ti, lh
I
I
as-polished 0.0 0.5 1.0 Potential / V vs SCE
Ni
2.0
1.5
Fig. 4. Cationic fraction of titanium in the surface film formed on the amorphous Ni-32Ti and Ni-62Ti alloys after potentiostatic polarization for 1, 2 and 4 h in deaerated 1 M HCI at 30 °C. The composition of the air-formed film after mechanical polishing is also shown for comparison.
Ti 1.0
+ 00 oO o= ~"
i
O [] z~ 0_
i
i
Ni-32Ti, lh Ni-32Ti, 2h Ni-32Ti, 4h Ni-62Ti, lh A
Ecorr " ~ E ~= 0.2 0.0
Ni
~
o ~ ' ~ I
I
i
a i = - 6 ~
0.6 0.4
i
deaerated 1M HCl 30"C
~
u
O
~
E]
Ni-32Ti I
I
as-polished 0.0 0.5 1.0 PotenUal / V vs SCE
I
1.5
2.0
Fig. 5. Atomic fraction of titanium in the underlying alloy surface just below the passive film on the amorphous Ni-32Ti and Ni-62Ti alloys after potentiostatic polarization for 1, 2 and 4 h in deaerated 1 M HCI at 30 °C. The alloy composition just below the air-formed film after mechanical polishing in cyclohexane is also shown for comparison.
the polarization potential. The composition of the airformed film and the composition of the underlying alloy surface just below the air-formed film are also shown for comparison. The titanium content in the air-formed film is higher than that of the bulk alloy, while the titanium content in the underlying alloy surface just below the air-formed film is lower than that of the bulk alloy. This result indicates the preferential oxidation of titanium in air. After open circuit immersion in 1 M HC1, the titanium content in the surface film is slightly higher than that of the air-formed film and the titanium content in the underlying surface of the Ni-62Ti alloy is decreased by immersion. Accordingly, it can be considered that preferential oxidation of titanium occurs by immersion. The content of titanium in the surface film decreases gradually on raising the potential. The cationic fraction of titanium in the film on the Ni-62Ti alloy is higher than that on the Ni-32Ti alloy. It can be assumed that the higher cationic fraction of titanium in the film on the Ni-62Ti alloy is responsible for the lower anodic current density. The composition of the surface film after prolonged anodic polarization is almost the same as that after polarization for 1 h. In the underlying alloy surface, the titanium content increases linearly with an increase in the potential. At relatively low potentials, nickel is concentrated in the underlying alloy surface, while at relatively high potentials the titanium content in the underlying alloy surface is slightly higher than that of the bulk alloy. The composition of the underlying alloy surface of the Ni-32Ti alloy after prolonged polarization at 0 V (SCE) becomes almost the same as the bulk alloy composition, but at + 1 and + 1.25 V (SCE) the composition is not changed by polarization for 2 or 4 h from that observed after polarization for 1 h. Accordingly the titanium-deficient underlying alloy surface is temporarily formed at low potentials and the steady composition of the underlying alloy surface at low potentials should be the same as the bulk alloy composition. The composition of the elements dissolved in 1 M HC1 during potentiostatic polarization at 0 and + 1 V (SCE) for 2 and 4 h for the amorphous Ni-32Ti alloy was measured by means of ICP analysis. The composition of elements dissolved was almost the same as the bulk alloy composition. It follows that there is no selectivity of elements for dissolution as a whole under the above conditions.
4. Discussion
The X-ray photoelectron spectroscopy (XPS) study revealed that the surface film tends to thicken linearly
E. A kiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni- Ti
with the polarization potential. The growth rate for the Ni-32Ti and Ni-62Ti alloys are the same as each other. Accordingly, the semiconducting properties of the passive film should be similar to each other. However, the surface film of the Ni-62Ti alloy is thicker than that of the Ni-32Ti alloy at each potential. The anodic current density of the Ni-62Ti alloy with a thicker film is lower than that of the Ni-32Ti alloy with a thinner film. In other words, the dissolution current at a fixed potential in the passive region decreases with decreasing film thickness. It seems that the observed film thickness does not necessarily indicate the thickness of the film grown anodically but the sum of the thickness of the film grown anodically and the thickness of the air-formed film components, i.e. mostly oxidized titanium without being dissolved by immersion or polarization in 1 M HC1. It can therefore be said that for these alloys some constituents of the airformed film act as a portion of barrier film along with the anodized film. The XPS study also revealed the enrichment of titanium in the surface film compared with the bulk alloy composition. The enrichment of titanium in the surface film in 1 M HCI is larger than that in 0.5 M H2SO 4 [10]. In other words, nickel is more difficult to use to constitute the passive film on the Ni-Ti alloys in hydrochloric acid solution in comparison with that formed in sulphuric acid solution. The open circuit immersion and polarization at low potentials give rise to an initial nickel enrichment in the underlying alloy surface as a result of preferential oxidation of titanium. However, when nickel is dissolved at high potentials a titanium enrichment occurs both in the film and in the underlying alloy surface. The thickness of the amorphous Ni-32Ti alloy dissolved which were estimated from the ICP analysis after potentiostatic polarization at + 1 V (SCE) for 2 h and 4 h were approximately 10 nm and 18 nm respectively. Furthermore, the IPC analysis of dissolved elements revealed no selective dissolution. Nevertheless, the enrichment of titanium in the underlying alloy surface was observed for the specimens after polarization at 1 V (SCE). In contrast, the thickening of titanium-enriched film by potentiostatic polarization at 1 V (SCE) for the alloys covered with the air-formed film is only about 2 nm as shown in Fig. 3. Since the density of the film is about half as large as that of the alloy, the titanium enrichment in the film should not affect the ratio of titanium to nickel obtained by the solution analysis. In addition, the titanium-enriched underlying surface should be quite thin so as to have no influence on the composition dissolved. It is interesting that the ratio of nickel to titanium dissolved from the Ni-Ti alloys under anodic polariza-
1131
tion through the titanium-enriched passive film is always almost the same as that in the bulk alloy composition. The enrichment of nickel or titanium in the underlying alloy surface is interpreted in terms of the difference in the oxidation rates of nickel and titanium depending upon the polarization potential. At low potentials titanium is mainly oxidized while at high potentials the rate of oxidation of nickel is higher than that of titanium at high potentials. Nevertheless, dissolution occurs following the bulk alloy composition because the element with a higher dissolution rate becomes deficient in the surface region. Accordingly, dissolution of metals following the bulk alloy composition at high potentials occurs through the titaniumenriched passive film and underlying alloy surface.
5. Conclusions Amorphous Ni-Ti alloys were prepared by d.c. magnetron sputtering. The composition range of single amorphous phase formation was about 30-70 at.%. The corrosion behaviour of these alloys in 1 M HC! was investigated in relation to the composition of the passive film and the underlying alloy surface just below the passive film. The amorphous Ni-Ti alloys were spontaneously passive in 1 M HC1. The passive film was enriched in titanium. The underlying alloy surface was initially enriched in nickel at relatively low polarization potentials, but was enriched in titanium at relatively high potentials. This fact is interpreted in terms of the difference in the oxidation rates of nickel and titanium depending on the polarization potential. That is to say, although the underlying alloy surface is enriched in either nickel or titanium, dissolution occurs following the bulk alloy composition.
Acknowledgment The present work is supported in part by Grant-inAid for Scientific Research (A) 03403012 from the Ministry of Education, Science and Culture, Japan.
References 1 M. Naka, K. Asami, K. Hashimoto and T. Masumoto, Proc, 4th Int. Conf. on Titanium, Metallurgical Society of AIME, Warrendale, PA, 1981, p. 2677. 2 A. Kawashima, K. Shimamura, S. Chiba, T. Matsunaga. K. Asami and K. Hashimoto, Proc. Asian-Pacific Corrosion Control Conf., Vol. 2, Tokyo, 1985, p. 1042. 3 K. Shimamura, A. Kawashima, A. Asami and K. Hashimoto, Sci. Rep. Res. Inst. Tohoku Univ., Ser. A, 33(1986) 196.
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E. Akiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni-Ti
4 A. Mitsuhashi, K. Asami, A. Kawashima and K. Hashimoto, Corros. Sei., 27(1987) 957. 5 A. Mitsuhashi, K. Asami, A. Kawashima and K. Hashimoto, in R. B. Diegle and K. Hashimoto (eds.), Proc. Symp. on Corrosion, Electrochemistry and Catalysis of Metallic Glasses, Electrochemical Society, Pennington, NJ, 1988, p. 191. 6 H. Yoshioka, A. Kawashima, K. Asami and K. Hashimoto, in R. B. Diegle and K. Hashimoto (eds.), Proc. Symp. on Corro-
7 8 9 10
sion, Electrochemistry and Catalysis of Metallic Glasses, Electrochemical Society, Pennington, NJ, 1988, p. 242. K. Asami, S.-C. Chen, H. Habazaki, A. Kawashima and K. Hashimoto, Corros. Sci., 31 (1990) 727. K. Asami, K. Hashimoto and S. Shimodaira, Corros. Sci., 17 (1977) 713. A. Aladjem, J. Mater. Sci., 8(1973)688. K. Asami, S.-C. Chen, H. Habazaki, A. Kawashima and K. Hashimoto, Corros. Sci., 31 (1990) 727.
1128
The corrosion behaviour of sputter-deposited amorphous Ni-Ti alloys in 1MHC1 Eiji Akiyama, Hiroki Habazaki, Asahi Kawashima, Katsuhiko Asami and Koji Hashimoto Institute for Materials Research, Tohoku University, Sendai 980 (Japan) Abstract Amorphous Ni-Ti alloys were prepared by d.c. magnetron sputtering. The corrosion behaviour of amorphous Ni-Ti alloys in 1 M HCI was investigated by electrochemical measurements. Particular attention was paid to the compositions of the surface film and underlying alloy surface just below the surface film. The amorphous Ni-Ti alloys were spontaneously passive. An X-ray photoelectron spectroscopy analysis revealed that the thickness of the surface film was increased linearly with an increase in the polarization potential. The growth rates for amorphous Ni-32Ti and Ni-62Ti alloys were the same as each other and slightly lower than the growth rate of sputter-deposited titanium. The passive film formed on the sputter-deposited amorphous Ni-Ti alloys was enriched in titanium cation. In the underlying alloy surface just below the surface film after polarization for 1 h, the titanium content increased linearly with the increase in the potential. At relatively low potentials, nickel was concentrated in the underlying alloy surface, while at relatively high potentials the content of titanium was higher than that in the bulk alloy.
1. Introduction
It has been found that the corrosion resistance of melt-spun amorphous Ni-Ti alloys is higher than that of the crystalline alloys [1], and that melt-spun amorphous Ni-Ta and Ni-Nb alloys have a high corrosion resistance even in hot concentrated HNO 3 [2], HC1 [2, 3] and H 3 P O 4 [4, 5].
However, the practical applications of the melt-spun amorphous alloys are limited by their shape. By contrast, sputer deposition is suitable for the preparation of the amorphous alloys with a wide surface area on conventional crystalline bulk metals. It is expected that the amorphous Ni-valve metal alloys prepared by sputtering also have high corrosion resistance. The present work aims to prepare corrosionresistant Ni-Ti alloys by a d.c. magnetron sputter deposition method and to clarify their corrosion and anodic polarization behaviour in 1 M HC1 solution. For a better understanding of the corrosion resistance of the alloys, their surface structures were analysed. Particular attention was paid to the compositions of the surface film and underlying alloy surface just below the surface film.
2. Experimental method
D.c. magnetron sputtering was carried out for preparation of Ni-Ti alloys. The target was composed of a 99.95% pure titanium disc of 100 mm diameter 0921-5093/94/$7.00 SSD10921-5093( 9 3 )05640-B
and 6 mm thickness, on the sputter erosion region of which 99.95% pure nickel discs of 20 mm diameter and 10 mm diameter were placed. The composition of the sputter deposits was changed by changing the numbers of nickel discs on the titanium disc. Glass plates were used as substrates. The sputtering apparatus and conditions used were the same as those described elsewhere [5, 6]. For the purpose of homogenization of the sputter deposits, the watercooled substrates were revolved around a central axis of the sputtering chamber, in addition to revolution of the substrates themselves around their own axes. After the target and substrates were installed in the sputtering machine the vacuum chamber was evacuated to about 1 x 10-6 Torr. After presputtering of the target for 5-10 min, sputtering was carried out at (4-8) x 10-4 Torr of argon gas. The composition of the alloys prepared was determined by inductively coupled plasma emission spectrometry (ICP) after the alloys were dissolved into 6 M HC1 solution. Their structure was identified by X-ray diffraction with Cu K a radiation. Potentiostatic polarization curves were measured in a deaerated 1 M HCI solution at 30 °C. In order to obtain reproducible results removal of the oxide film formed on as-sputtered alloys by air exposure was necessary, and hence, before polarization or immersion, the specimens were polished mechanically with silicon carbide paper up to grit 1500 in cyclohexane and dried in air. A platinum electrode and a saturated calomel elecrode (SCE) were used as counterelectrode © 1994 - Elsevier Sequoia. All rights reserved
E. Akiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni-Ti
and reference electrode respectively. Potentials hereafter are referred to the SCE. After potentiostatic polarization, X-ray photoelectron spectra were measured by means of a Shimadzu-850 photoelectron spectrometer with Mg K a radiation. For quantitative analysis of the alloy surface, the photoionization cross-section of the Ni 2p3/2 electrons relative to the O ls electrons including the satellite of the Ni 2p3/2 spectrum was determined by using the standard substances Ni(OH)2 and Ni203 as 2.315. The photoionization cross-section of the Ti 2p3/2 electrons relative to O ls electrons used was 1.277 [7]. The thickness and composition of the surface film and the composition of the underlying alloy surface were quantitatively determined by a previously proposed method using integrated intensities of photoelectrons [8]. A solution analysis was performed by means of ICE The concentrations of nickel and titanium dissolved were measured at 231.604 nm and 334.941 nm wavelengths respectively.
range. The broad peak at about 23 ° arises from the glass substrate. When the alloy titanium content increases, the continuous shift of the first halo position was observed from that close to the 111 reflection of the f.c.c, nickel phase to that close to the 011 reflection of the h.c.p, titanium phase. This indicates that the alloy consists of a single amorphous phase. Figure 2 shows the potentiostatic polarization curves for the amorphous Ni-32Ti and Ni-62Ti alloys. The amorphous Ni-Ti alloys are spontaneously passive and their anodic current densities are almost constant in the passive region. An abrupt increase of the current density is observed at about 1.5 V. This current increase is attributable to oxygen evolution. The current density of the Ni-32Ti alloy is higher than that of the Ni-62Ti alloy. Accordingly the corrosion resistance of the Ni-62Ti alloy is higher than that of the Ni-32Ti alloy. Figure 3 shows the change in the thickness of the surface film formed on the amorphous Ni-32Ti and Ni-62Ti alloys and sputter-deposited titanium as a
10 3
3. Experimental results
l
101
I
//
Ti sputtered
=- 10 0 e,-
[
l
o;c
'E
a
I
deaerated 1 M HCI
¢q
Figure 1 shows X-ray diffraction patterns of sputterdeposited Ni-Ti alloys. Alloy compositions hereafter are all denoted in atomic percentages. When the titanium content is very low, such as in Ni-16Ti alloy, reflections close to the 111,200 and 220 reflections of the f.c.c, nickel phase are observed. The alloys containing 32 at.% or more of Ti show the halo pattern typical for the amorphous structure over a wide composition
1129
//
....\
10 1
//
Ni-32Ti
:~
f
= 10-2
o 10 -3
I
-1.0
I
-0.5
0.0
0.5
1.0
1.5
2.0
P o t e n t i a l / V v s SCE
I
I
I
I
I
Fig. 2. Potentiostatic polarization curves of sputter-deposited amorphous Ni-32Ti and N i - 6 2 T i alloys in deaerated 1 M HC1 at 30 °C.
I
70
I
I
I
deaerated 1M HCI o
8 C
50
._o .c= I-- 40 IE 30
as-polished
~
-1.0
Ni-7,1Ti
I
Ni-62Ti
2Ti
qr 20
I
___~ -~//'~~
" 13
30 C, l h
,~ 6O
Ti sputtered
Ni-62Ti : 1 6 ~ v '
J
I
I
I
-0.5
0.0
0.5
1.0
I
1.5
2.0
Potential / V vs SCE
20
I
I
I
i
I
I
30
40
50
60
70
80
90
2 0 / degree Fig. 1. X-ray diffraction patterns of sputter-deposited Ni-Ti alloys.
Fig. 3. Thickness of surface films formed on the amorphous Ni-32Ti and Ni-62Ti alloys and sputter-deposited titanium in 1 M HC1 at 30 °C. The thickness of the air-formed film after mechanical polishing in cyclohexane is also shown for comparison.
1130
E. Akiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni-Ti
function of polarization potential. This figure clearly reveals that anodic polarization leads to film thickening with the polarization potential. The apparent growth rates calculated from the slope of the straight lines for the amorphous Ni-32Ti and Ni-62Ti alloys are both approximately 16 A V- 1. The growth rate for sputterdeposited titanium is 21 A W-1, which is almost the same as the anodizing ratio reported for titanium [9]. The growth rate for the amorphous Ni-Ti alloys is smaller than that for sputtered titanium, perhaps because the surface film contains nickel ions. Figures 4 and 5 show the cationic fraction of titanium in the passive film and the atomic fraction of titanium in the underlying alloy surface just below the passive film for the amorphous Ni-32Ti and Ni-62Ti alloys after anodic polarization for 1 h as a function of Ti
E 1.0 i,
= E 0.8 _= e= ¢0
~- 0.6
"6 tO
0.4
O [] z~ •
o
.,= U.
_o 0.2 e. o
o 0.0
deaerated 1M HCI 30°C I
I
I
Ni-32Ti, lh Ni-32Ti, 2h Ni-32Ti, 4h Ni-62Ti, lh
I
I
as-polished 0.0 0.5 1.0 Potential / V vs SCE
Ni
2.0
1.5
Fig. 4. Cationic fraction of titanium in the surface film formed on the amorphous Ni-32Ti and Ni-62Ti alloys after potentiostatic polarization for 1, 2 and 4 h in deaerated 1 M HCI at 30 °C. The composition of the air-formed film after mechanical polishing is also shown for comparison.
Ti 1.0
+ 00 oO o= ~"
i
O [] z~ 0_
i
i
Ni-32Ti, lh Ni-32Ti, 2h Ni-32Ti, 4h Ni-62Ti, lh A
Ecorr " ~ E ~= 0.2 0.0
Ni
~
o ~ ' ~ I
I
i
a i = - 6 ~
0.6 0.4
i
deaerated 1M HCl 30"C
~
u
O
~
E]
Ni-32Ti I
I
as-polished 0.0 0.5 1.0 PotenUal / V vs SCE
I
1.5
2.0
Fig. 5. Atomic fraction of titanium in the underlying alloy surface just below the passive film on the amorphous Ni-32Ti and Ni-62Ti alloys after potentiostatic polarization for 1, 2 and 4 h in deaerated 1 M HCI at 30 °C. The alloy composition just below the air-formed film after mechanical polishing in cyclohexane is also shown for comparison.
the polarization potential. The composition of the airformed film and the composition of the underlying alloy surface just below the air-formed film are also shown for comparison. The titanium content in the air-formed film is higher than that of the bulk alloy, while the titanium content in the underlying alloy surface just below the air-formed film is lower than that of the bulk alloy. This result indicates the preferential oxidation of titanium in air. After open circuit immersion in 1 M HC1, the titanium content in the surface film is slightly higher than that of the air-formed film and the titanium content in the underlying surface of the Ni-62Ti alloy is decreased by immersion. Accordingly, it can be considered that preferential oxidation of titanium occurs by immersion. The content of titanium in the surface film decreases gradually on raising the potential. The cationic fraction of titanium in the film on the Ni-62Ti alloy is higher than that on the Ni-32Ti alloy. It can be assumed that the higher cationic fraction of titanium in the film on the Ni-62Ti alloy is responsible for the lower anodic current density. The composition of the surface film after prolonged anodic polarization is almost the same as that after polarization for 1 h. In the underlying alloy surface, the titanium content increases linearly with an increase in the potential. At relatively low potentials, nickel is concentrated in the underlying alloy surface, while at relatively high potentials the titanium content in the underlying alloy surface is slightly higher than that of the bulk alloy. The composition of the underlying alloy surface of the Ni-32Ti alloy after prolonged polarization at 0 V (SCE) becomes almost the same as the bulk alloy composition, but at + 1 and + 1.25 V (SCE) the composition is not changed by polarization for 2 or 4 h from that observed after polarization for 1 h. Accordingly the titanium-deficient underlying alloy surface is temporarily formed at low potentials and the steady composition of the underlying alloy surface at low potentials should be the same as the bulk alloy composition. The composition of the elements dissolved in 1 M HC1 during potentiostatic polarization at 0 and + 1 V (SCE) for 2 and 4 h for the amorphous Ni-32Ti alloy was measured by means of ICP analysis. The composition of elements dissolved was almost the same as the bulk alloy composition. It follows that there is no selectivity of elements for dissolution as a whole under the above conditions.
4. Discussion
The X-ray photoelectron spectroscopy (XPS) study revealed that the surface film tends to thicken linearly
E. A kiyama et al.
/
Corrosion behaviour of sputter-deposited amorphous Ni- Ti
with the polarization potential. The growth rate for the Ni-32Ti and Ni-62Ti alloys are the same as each other. Accordingly, the semiconducting properties of the passive film should be similar to each other. However, the surface film of the Ni-62Ti alloy is thicker than that of the Ni-32Ti alloy at each potential. The anodic current density of the Ni-62Ti alloy with a thicker film is lower than that of the Ni-32Ti alloy with a thinner film. In other words, the dissolution current at a fixed potential in the passive region decreases with decreasing film thickness. It seems that the observed film thickness does not necessarily indicate the thickness of the film grown anodically but the sum of the thickness of the film grown anodically and the thickness of the air-formed film components, i.e. mostly oxidized titanium without being dissolved by immersion or polarization in 1 M HC1. It can therefore be said that for these alloys some constituents of the airformed film act as a portion of barrier film along with the anodized film. The XPS study also revealed the enrichment of titanium in the surface film compared with the bulk alloy composition. The enrichment of titanium in the surface film in 1 M HCI is larger than that in 0.5 M H2SO 4 [10]. In other words, nickel is more difficult to use to constitute the passive film on the Ni-Ti alloys in hydrochloric acid solution in comparison with that formed in sulphuric acid solution. The open circuit immersion and polarization at low potentials give rise to an initial nickel enrichment in the underlying alloy surface as a result of preferential oxidation of titanium. However, when nickel is dissolved at high potentials a titanium enrichment occurs both in the film and in the underlying alloy surface. The thickness of the amorphous Ni-32Ti alloy dissolved which were estimated from the ICP analysis after potentiostatic polarization at + 1 V (SCE) for 2 h and 4 h were approximately 10 nm and 18 nm respectively. Furthermore, the IPC analysis of dissolved elements revealed no selective dissolution. Nevertheless, the enrichment of titanium in the underlying alloy surface was observed for the specimens after polarization at 1 V (SCE). In contrast, the thickening of titanium-enriched film by potentiostatic polarization at 1 V (SCE) for the alloys covered with the air-formed film is only about 2 nm as shown in Fig. 3. Since the density of the film is about half as large as that of the alloy, the titanium enrichment in the film should not affect the ratio of titanium to nickel obtained by the solution analysis. In addition, the titanium-enriched underlying surface should be quite thin so as to have no influence on the composition dissolved. It is interesting that the ratio of nickel to titanium dissolved from the Ni-Ti alloys under anodic polariza-
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tion through the titanium-enriched passive film is always almost the same as that in the bulk alloy composition. The enrichment of nickel or titanium in the underlying alloy surface is interpreted in terms of the difference in the oxidation rates of nickel and titanium depending upon the polarization potential. At low potentials titanium is mainly oxidized while at high potentials the rate of oxidation of nickel is higher than that of titanium at high potentials. Nevertheless, dissolution occurs following the bulk alloy composition because the element with a higher dissolution rate becomes deficient in the surface region. Accordingly, dissolution of metals following the bulk alloy composition at high potentials occurs through the titaniumenriched passive film and underlying alloy surface.
5. Conclusions Amorphous Ni-Ti alloys were prepared by d.c. magnetron sputtering. The composition range of single amorphous phase formation was about 30-70 at.%. The corrosion behaviour of these alloys in 1 M HC! was investigated in relation to the composition of the passive film and the underlying alloy surface just below the passive film. The amorphous Ni-Ti alloys were spontaneously passive in 1 M HC1. The passive film was enriched in titanium. The underlying alloy surface was initially enriched in nickel at relatively low polarization potentials, but was enriched in titanium at relatively high potentials. This fact is interpreted in terms of the difference in the oxidation rates of nickel and titanium depending on the polarization potential. That is to say, although the underlying alloy surface is enriched in either nickel or titanium, dissolution occurs following the bulk alloy composition.
Acknowledgment The present work is supported in part by Grant-inAid for Scientific Research (A) 03403012 from the Ministry of Education, Science and Culture, Japan.
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