Oxidation of TiNi surface with hyperthermal oxygen molecular beams

Applied Surface Science 257 (2011) 4257–4263

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Oxidation of TiNi surface with hyperthermal oxygen molecular beams Michio Okada a,b,∗ , Makoto Souwa b , Toshio Kasai b , Yuden Teraoka c a

Renovation Center of Instruments for Science Education and Technology, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan Department of Chemistry, Graduate School of Science, Osaka University, 1-1 Mchikaneyama-cho, Toyonaka, Osaka 560-0043, Japan c Synchrotron Radiation Research Center, Japan Atomic Energy Agency, 1-1-1 Kouto, Sayo-cho, Sayo, Hyogo 679-5148, Japan b

a r t i c l e

i n f o

Article history: Received 24 November 2010 Received in revised form 6 December 2010 Accepted 6 December 2010 Available online 22 December 2010 Keywords: TiNi alloy Hyperthermal molecular beam Oxidation

a b s t r a c t We report results of our detailed studies on the initial oxidation process of TiNi with a 2 eV hyperthermal oxygen molecular beam (HOMB) and thermal O2 in the backfilling. The oxidation processes are monitored by X-ray photoemission spectroscopy (XPS) measurements in conjunction with synchrotron radiation (SR). In the early stage of oxidation, the precursor mediated dissociative adsorption is the dominant reaction mechanism. In the oxide formation process at higher O coverage, HOMB has the advantage in the dissociation process of O2 molecule and can grow TiO2 layers with the underlying TiOx -rich and/or Ni-rich layers. We succeeded in fabricating thick Ni-free TiO2 layer, possibly blue colored rutile TiO2 , combining HOMB and surface annealing. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Ti–Ni alloys with nearly equiatomic nickel to titanium ratios exhibit attractive properties, such as superelasticity with a large recoverable strain shape memory effect, excellent corrosion resistance and biocompatibility. TiNi alloy is one of the most useful shape memory alloys (SMAs) because of a shape memory effect up to 8% strain [1]. In recent years, it has been used in industrial applications as well as many biomedical applications [2,3]. Various medical applications were expected on TiNi SAM [4,5] from the early stages, although Ni may induce medical problems [6]. There have been a lot of studies on the oxidation of Ti–Ni alloys at various atmosphere and temperatures, and from various points of views [7–24]. In some cases, a TiO2 rich and Ni depleted zone is formed below the surface mixed phase of Ti and Ni oxides [9,10,17,25]. In other cases, almost Ni free TiO2 is formed on the surface region [14,25]. For good biocompatibility, removing the toxicity hazard, the release of Ni ions must be limited [25–29]. For that purpose, a Ni-free titanium oxide layer must be fabricated as a protective layer on the Ti–Ni alloy surface. Hyperthermal O2 molecular beam (HOMB) is an efficient tool fabricating oxide layers on various metal surfaces [30–32] even at room temperature (RT) and lower. On a Cu3 Au surface [31], HOMB induces Cu segregation on the surface and produces protective Cu–O layer with Au in the second layer even in the conditions where thermal O2 dose produces no effective oxidation. Efficient oxida-

∗ Corresponding author. Tel.: +81 6 6879 4783; fax: +81 6 6879 4783. E-mail address: [email protected] (M. Okada). 0169-4332/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2010.12.033

tion of TiNi requires high surface temperature in various chemical environments [10–12,17,18]. It is expected that HOMB can improve the oxide formation due to its kinetic energy dissipation during the collision and reactions. Herein we report results of our detailed studies on the initial oxidation process of TiNi with a 2 eV HOMB and thermal O2 in the backfilling. From the O-uptake curves, which were determined from a series of O-1s X-ray photoemission spectroscopy (XPS) measurements in conjunction with synchrotron radiation (SR), it was found that the dissociative adsorption of O2 is less efficient at low O coverage for the 2 eV HOMB dose than thermal O2 in the backfilling. The SR-XPS measurements suggest the dissociative adsorption of O2 occurs in accompany with the Ti segregation on the surface in agreement with the documented works in the vacuum [7,9]. The growth of Ti2 O for 2 eV HOMB suggests that the diffusion of Ti and Ni atoms also contribute dominantly to the oxidation process. 2. Experiments All experiments were performed using the surface reaction analysis apparatus (SUREAC 2000) constructed at BL23SU in SPring-8 [33]. Briefly, the surface reaction analysis chamber is equipped with an electron energy analyzer (OMICRON EA125) and a Mg/Al-Ka twin-anode X-ray source (OMICRON DAR400). A quadrupole mass spectrometer, which was used to analyze the molecular species in the HOMB, was located on the opposite side of the HOMB source. The base pressure of the surface reaction chamber was below 1 × 10−8 Pa. The TiNi alloy sample (7 mm × 7 mm × 1 mm, purity of 99.9%, KOJUNDO Chem. Lab. Co.) was polished mechanically and washed with acetone, ethanol, and de-ionized water in ultrasonic

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bath. According to X-ray fluorescence analysis, the atomic concentration of bulk Ti and Ni was 48.9% and 51.1%. The TiNi sample was cleaned by repeatedly sputtering with 1-keV Ar+ for an hour until no impurities except for O atoms were detected by SR-XPS. In order to reduce the residual O atoms on the surface as low as possible, only sputtering was performed with no annealing. Annealing procedure make surface O atoms increase during the cooling of the TiNi sample due to its high reactivity with the residual gas. Preferential sputtering of Ti [9,17] occurred on this surface and the number ratio of surface atoms Ni/Ti was 1.12 after the cleaning procedures. Thus, we prepared for the initial surface so carefully as to reproduce the XPS spectra. The incident 2 eV HOMB was produced by the expansion of the mixture gas of He (99.9%) and O2 (1%) at the nozzle temperature of 1400 K. The O2 flux at the sample position is estimated experimentally before and/or after each measurement, following the estimation method described in details in Ref. [33]. The typical flux density of O2 molecules at the sample surface is 2 × 1015 molecules cm−2 s−1 for the 2 eV HOMB. In the thermal O2 dose, the sample was exposed to 1.33 × 10−5 Pa O2 by backfilling the chamber. The typical exposure unit of 1 Langmuir (L) corresponds to 1.33 × 10−4 Pa s and 3.58 × 1014 O2 molecules cm−2 at 300 K. After irradiating the TiNi surface with a proper amount of HOMB in the surface normal or thermal O2 , high-resolution SRXPS spectra were measured at 0◦ and 70◦ from the surface normal using a monochromatic SR beam with photon energies of 1156 and 1158 eV. Dosing oxygen was carried out at the sample temperature of 300 or 673 K. All XPS measurements were performed at a sample temperature of 300 K. Transmission electron microscope (TEM) observations and energy dispersive X-ray fluorescence analysis were performed for the sufficiently oxidized TiNi samples. TEM observations were performed at electron beam energy of 200 keV with JEM-2010F in KOBELCO Research Institute, Inc. Energy dispersive X-ray (EDX) spectrometry measurements were performed with the beam radius of ∼1 nm. 3. Results and discussions Fig. 1(a)–(c) shows the HOMB exposure dependence of O 1s, Ti 2p and Ni 2p XPS spectra at room temperature (RT), respectively. The following are the summary of the features in Fig. 1(a)–(c). (1) With increasing HOMB exposure, the peak intensity increases and the peak position of O 1s shifts from 529.9 eV to 530.5 eV by 0.6 eV. (2) The intensity of Ti 2p peaks at 453.8 and 459.9 eV appearing on the clean surface decrease with increasing HOMB exposure. Alternatively, the Ti 2p peaks at 458.8 and 464.6 eV grow, corresponding to the oxidation of Ti. (3) The intensity of Ni 2p peaks at 852.8 and 869.9 eV appearing on the clean surface decrease with increasing HOMB exposure. No Ni oxide peaks appeared even after prolonged oxidation. These results suggest that the preferential oxidation of Ti occurs with the Ti segregation on the surface and as a result, Ni moves to bulk below the Ti/Ti oxide layers. We separate the Ti2p peak into each oxide component according to oxidation numbers, as shown in Fig. 2. The peak fitting was performed using UNIFIT2002 software after subtracting the Shirley background. The peaks were fitted with Voigt function. The line shape of each component is fixed and common for all HOMB exposures. Peak positions were set to the same ones as previously reported for TiNi [14–16]: TiO2 (Ti4+ , 2p1/2 : 464.6 eV, 2p3/2 : 458.8 eV), Ti2 O3 (Ti3+ , 2p1/2 : 463.1 eV, 2p3/2 : 457.0 eV), TiO (Ti2+ , 2p1/2 : 461.5 eV, 2p3/2 : 455.3 eV), and Ti (Ti0 (TiA 0 ), 2p1/2 : 460.5 eV,

Fig. 1. 2 eV HOMB exposure dependence of (a) O 1s, (b) Ti 2p, and (c) Ni 2p measured in the surface sensitive geometry on TiNi at RT.

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Fig. 2. Peak separation of Ti 2p into each oxidation state of Ti for (a) the surface-normal detection and (b) the surface sensitive detection, in the 2 eV HOMB exposure on TiNi at RT.

2p3/2 : 454.5 eV). For the best fitting, additional component (TiB 0 , 2p1/2 : 461.0 eV, 2p3/2 : 455.0 eV) corresponding to Ti0 is required at the higher binding energy compared to the previously reported one TiA 0 . The peak position of TiB 0 is closer to the Ti0 binding energy measured on Ti (0 0 0 1) [34] than that of TiA 0 . Thus, we can tentatively assign TiA 0 and TiB 0 to metallic Ti in the Ni-rich and the Ti-rich sites, respectively. This is consistent with the faster decay of TiB 0 during the oxidation caused by the preferential oxidation of Ti. Ti surrounded by Ti is oxidized more easily than Ti surrounded by Ni. It should be noted that the free energy of formation G (298 K) for NiO, TiO and TiO2 is −211.7, −495, and −889.5 kJ mol−1 , respectively [35]. Ti being the more reactive element, segregates to the surface and is oxidized there forming TiO2 [7] with a Ni enriched alloy layer underneath. Fig. 2 demonstrates that HOMB induces the Ti oxide formation, especially TiO2 formation. In the mean HOMB exposure of 1.1 × 1016 molecules cm−2 , all oxidation states are observed with comparable intensities. This result suggests the sequential oxidation from Ti to TiO2 in the early stage of oxidation. Further increase of HOMB dose induces the stable TiO2 formation. Almost 90% of metal components in the surface region form TiO2 at 5.23 × 1018 molecules cm−2 from the surface sensitive detection (∼1.2 nm from the surface), while there are still other oxidation states, especially Ti2 O3 , in the further deeper layer from the surfacenormal detection (∼3 nm from the surface). At RT, oxide thickness is limited (vide infra) probably due to the low diffusion of Ti and Ni. Thus, the lower oxide states are almost completely oxidized to TiO2 by the enough available O atoms. The layer composition may vary with the phase stability depending on the concentration of each element [10]. Fig. 3 shows the TEM image of the 2 eV HOMB oxidized surface. The TiNi sample was taken out in the air after the 5.6 × 1019 molecules cm−2 HOMB exposure and then, the TEM image of the sample was measured. The XPS spectra are quite similar to those for the highest exposure in Fig. 1. Three regions of A, B and C can be separated; A: 0–2 nm from the surface, B: 2–7 nm from the surface, C: >7 nm form the surface. From the EDX measurements, A and B correspond to the oxide layers. The 90% of metal atoms in the region <1.2 nm of A form TiO2 , as determined from the surface sensitive XPS detection, while the Ni/Ti atomic ratio in region A is 0.15 determined from the EDX measurements. These

results suggest that Ti segregates to the surface and forms stable TiO2 (see the above-mentioned formation free energy). The region B contains the considerable amounts of Ti2 O3 and the other states in addition to TiO2 , as determined from the XPS surface-normal detection. The Ni/Ti atomic ratio in region B is 0.89 obtained from the EDX measurements. In region C, the Ni/Ti atomic ratio is 0.53, suggesting the Ti segregation from the far bulk (Ni/Ti ∼ 1). The observation of the Ni rich region beneath the Ti oxide layer agrees with the previous reports [7,8] in high temperature oxidation of TiNi. The composition of Ti-oxide states depends on the depth from the surface [10,14]. In the oxidation of TiNi, the metal-element diffusion is important than the O diffusion [11]. From the diffusivity measurements of metal in Ti [36], Ni in Ti diffuses 106 times faster than Ti in Ni. Furthermore, the Ni diffusivity is high in the defective oxide states TiOx (x < 2) [8], and the Ti diffusivity is also high in TiO2 [11]. Thus, during the oxidation, Ni atoms diffuse into bulk and at the

Fig. 3. TEM image of the oxidized TiNi for the 2 eV HOMB exposure of 5.6 × 1019 O2 cm−2 at RT.

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Fig. 4. Incident energy dependence of uptake curve for the O2 exposure on TiNi at RT, determined by the O 1s evolution.

same time the Ti segregation occurs, via the defective oxide sites TiOx (x < 2). It should be noted that the efficiency of TiO2 -layer formation is lower on TiAl [37] where the diffusivity of Al in Ti is as low as that of Ti in Ti [36]. Next, we compare the oxidation with 2 eV HOMB and thermal O2 dose in backfilling (∼25 meV) in order to discuss the effect of the translational energy of incident O2 molecules. Fig. 4 shows the incident energy dependence of the uptake curve obtained from the integration of the O 1s XPS peaks. At low exposures below 3 × 1016 molecules cm−2 , the oxidation efficiency in the thermal O2 dose is higher than that in the HOMB incidence. This result suggests that the precursor-mediated dissociative adsorption of O2 molecules is dominant in the initial stage of oxidation. Molecularly adsorbed O2 precursor is mobile and dissociates at the lower activation sites, possibly Ti rich sites. The trapping probability in such precursor states is higher for thermal O2 than for HOMB, because HOMB must suffer much larger energy dissipation for the trapping. Further oxidation at higher exposures above 3 × 1016 molecules cm−2 requires additional energy for the dissociation, due to the higher activation barrier for the dissociation caused by the repulsive interaction from the preadsorbed O atoms. Moreover, HOMB collision induced absorption of O atoms on the surface creates new dissociation sites [30]. Thus, the higher incident energy induces the more effective oxidation in the higher exposures in Fig. 4. Fig. 5(a) and (b) shows the evolution of Ti 2p and Ni 2p XPS peak intensity during the oxidation in Fig. 4. The intensity of Ti 2p increases with increasing oxidation dose. On the other hand, the intensity of Ni 2p decreases with increasing oxidation dose. More efficient Ti segregation and Ni diffusion at the low thermal O2 doses suggest that the Ti segregation and the Ni diffusion are not enhanced by the translational energy of incident molecules. Comparison of Fig. 5(a) and (b) with Fig. 4 suggests that the Ti and the Ni intensities seem to depend only on the O coverage. In order to confirm the origin of the evolution of Ti 2p and Ni 2p XPS peak intensity, we re-plot the data of Fig. 5(a) and (b) as a function of O 1s XPS intensity in Fig. 5(c) and (d), respectively. The Ti 2p and Ni 2p XPS intensity is scaled by the O coverage and then, the amounts of Ti and Ni in the surface region depend not on the incident energy but just on the O coverage. Thus, it is concluded that the incident energy of O2 molecules contributes to the efficiency of oxidation, the dissociation processes, and as a result of oxidation, the Ti seg-

regation and the Ni diffusion are induced. Similar conclusion was obtained from the surface-sensitive detection (not shown here). As discussed above, the Ni thermal diffusion plays an important role in the Ti segregation and is also enhanced by the gradient of chemical potential during the oxide formation in the surface region. Fig. 6 shows the oxidation-state dependence of Ti 2p peak evolution plotted as a function of O 1s XPS intensity. The TiO2 (Ti4+ ) component increases monotonically with increasing O 1s intensity, suggesting the most stable oxide of TiO2 . On the other hand, the Ti2 O3 component increases in the early sage of oxidation and then saturates. This is consistent with the TEM observation and the XPS result that TiO2 is formed in the surface region and Ti2 O3 is located mainly in the deeper layers, probably in the interface layers between TiO2 and TiNi metal. The oxide state of Ti2 O3 may be in the equilibrium state depending on the layer composition, during the oxidation. Ti2 O3 is also a defective state [8,9] where metal diffusion possibly occurs. It is interesting that the amount of Ti2 O3 is larger for the thermal O2 dose, in the surface-normal detection, than for HOMB incidence, while no differences were observed in the surface-sensitive detection. The Ti2 O3 state, which corresponds to the defective precursors for TiO2 formation, is formed on the surface at the first stage of oxidation and then is buried mainly in region B under the surface region of A in Fig. 3 during the oxidation at room temperature. Thus, the first stage of formation of Ti2 O3 is affected by the incident energy of O2 molecules. HOMB can produce the less defective surface [38], i.e., less Ti2 O3 during its energy dissipation. The buried interface B may keep the early stage of traces occurring on the topmost layer in the oxidation. In order to fabricate thicker TiO2 layers efficiently, we tried to combine the HOMB incidence and the thermal annealing of the surface. We expose 2-eV HOMB on the clean TiNi at the surface temperature of 673 K. Fig. 7(a) and (b) shows the exposure dependence of Ti 2p and O 1s XPS spectra, respectively. At 1.60 × 1019 molecules cm−2 , various oxide states are observed comparably, which is quite different from the oxidation at RT [24]. At RT, ∼5 × 1018 molecules cm−2 is enough to grow the TiO2 layer although the thickness is thin. The corresponding O 1s XPS spectrum shows the broad peak shape. Higher temperature enhances the diffusion of Ti atoms to the surface and Ni atoms into the bulk. Thus, the supply of Ti atoms from the bulk is enough to produce thicker TiO2 if the ample amounts of dissociated O atoms are available. The present experimental result suggests that the supply of O atoms from the dissociation of O2 molecules is insufficient for the TiO2 formation, causing a kinetic limit of TiO2 growth. At 8.39 × 1019 molecules cm−2 , TiO2 is grown accompanied with a small amount of Ti2 O3 . 90% of the surface region monitored by the XPS measurement is TiO2 . The corresponding O 1s XPS spectrum becomes sharper than at 1.60 × 1019 molecules cm−2 , suggesting the dominant TiO2 growth. Fig. 8 shows the TEM image of the 2 eV HOMB oxidized surface at 673 K. corresponding to the highest exposure of HOMB in Fig. 7. Four regions of A , B , C , D can be separated. A contains 90 nm TiO2 layers and B is also 40 nm Ti rich oxide layers. The crystalline features were observed in A . From the EDX measurements, the ratio Ni/Ti is 0.05 and 0.09, respectively. Although we cannot measure the XPS data for B , it is considered from the TEM image in Fig. 3 that B contains various oxide states. In region C below B , a small amount of O atoms was detected in the EDX measurements and thus, no considerable oxidation proceeds. The ratio Ni/Ti determined by the EDX measurements is 2.25. Thus, Ni atoms in the surface region diffuse into the region C and Ti atoms are crowded out into the region B . In the further deeper region D , no O atoms were detected by EDX and the ratio Ni/Ti is 1.09 which is almost bulk value. Fig. 8 also shows the photograph of the oxidized TiNi taken before the TEM measurements. The blue-colored part was shined by HOMB and contains thick TiO2 layers. Thus,

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Fig. 5. Incident energy dependence of (a) the Ti 2p and (b) the Ni 2p intensities for the O2 exposure on TiNi at RT. The evolutions of the Ti 2p and the Ni 2p intensities are re-plotted as a function of O 1s intensity in (c) and (d), respectively.

Fig. 6. The evolutions of the TiO2 and the Ti2 O3 intensities as a function of O 1s intensity for (a) the surface-normal detection and (b) the surface-sensitive detection of Ti 2p in the O2 exposure on TiNi at RT.

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Fig. 7. Evolutions of (a) Ti 2p and (b) O 1s measured in the surface-normal geometry for 2 eV HOMB exposure on TiNi at 673 K.

the blue-colored rutile TiO2 [10,13] may be grown by this method. Fig. 9 shows the Ni 2p XPS spectrum of this heavily oxidized surface, compared to that of the clean surface. After oxidation, no Ni was detected within our experimental accuracy. Thus, we could

produce perfectly Ni-free TiO2 protective layer on the surface. Combination of HOMB and annealing is an efficient tool for fabricating Ni-free TiO2 films on the surface and can be used for biocompatible material fabrication. 4. Summary We report results of our detailed studies on the initial oxidation process of TiNi with a 2 eV HOMB and thermal O2 in the backfilling. The dissociative adsorption of O2 is less efficient at low O coverage for the 2 eV HOMB dose than thermal O2 in the backfilling. In the early stage of oxidation, the precursor mediated dissociative adsorption is the dominant reaction mechanism. In the oxide formation process at higher O coverage, HOMB has the advantage in the dissociation process of O2 molecule and growing TiO2 layers with the underlying TiOx -rich and/or Ni-rich layers. Ti segregation

Fig. 8. TEM image of the oxidized TiNi corresponding to the highest exposure in Fig. 8. Inset shows the photograph of the corresponding TiNi sample.

Fig. 9. Exposure dependence of Ni 2p measured in the surface-normal geometry for 2 eV HOMB exposure on TiNi at 673 K.

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and Ni diffusion through Ti2 O3 defective sites process plays a key role in the oxide formation. We succeeded in fabricating thick Nifree TiO2 layer, possibly blue colored rutile TiO2 , combining HOMB and surface annealing. HOMB is a powerful tool for fabricating oxide thin films. Acknowledgements The experiments were performed using SUREAC 2000 in BL23SU at SPring-8 with the approval of the Japan Synchrotron Radiation Research Institute (JASRI) and Japan Atomic Energy Agency (JAEA). The authors are grateful to Dr. Y. Saitoh and Dr. A. Yoshigoe for their help in operating the monochromatic system at the beam line. A part of this work has been supported by New Energy and Industrial Technology Development Organization (NEDO) under “Advanced Fundamental Research Project on Hydrogen Storage Materials”. The Japanese Ministry of Education, Culture, Sports, Science and Technology is gratefully acknowledged for the Grant-in-Aid for Scientific Research (Nos. 20350005, 20360024, 22655005) that supported partially of this work. MO was also financially supported by PRESTO of JST, The Mitsubishi Foundation and The Sumitomo Foundation. References [1] W.J. Buehler, J.V. Gifrich, R.C. Wiley, J. Appl. Phys. 34 (1963) 1475. [2] B. Thierry, M. Tabrizian, C. Trepanier, O. Savadogo, L.H. Yahia, J. Biomed. Mater. Res. 51 (2000) 685. [3] D.J. Wever, A.G. Veldhuizen, J. de Vries, H.J. Busscher, D.R.A. Uges, J.R. van Horn, Biomaterials 19 (1998) 761. [4] J. Haasters, G.V. Salis-Solio, G. Bensmann, in: T.W. Duerig, K.L. Melton, D. Stöckel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloy, ButterworthHeinemann, UK, 1990, p. 426. [5] H. Onishi, Jpn. Soc. Artif. Organs 12 (1983) 862 (in Japanese). [6] S. Kawanishi, S. Inoue, S. Oikawa, N. Yamashita, S. Toyokuni, M. Kawanishi, K. Nishino, Free Radical Biol. Med. 31 (2001) 108. [7] P.H. McBreen, M. Polak, Surf. Sci. 179 (1979) 483. [8] J.P. Espinós, A. Fernández, A.R. González-Elipe, Surf. Sci. 295 (1993) 402. [9] C.-M. Chan, S. Trigwell, T. Duerig, Surf. Interface Anal. 15 (1990) 349.

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