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Electron beam induced oxidation of Ni3Al surfaces: electron flux effects
Surface Science 507–510 (2002) 486–491 www.elsevier.com/locate/susc
Electron beam induced oxidation of Ni3Al surfaces: electron flux effects S.A. Koch, G. Palasantzas, D.T.L. van Agterveld, J.Th.M. De Hosson
*
Department of Applied Physics, Materials Science Center and Netherlands Institute for Metals Research, University of Groningen, Nijenborgh 4, 9747 AG Groningen, Netherlands
Abstract Electron beam irradiation of polycrystalline boron doped Ni3 Al (at 300 K and under ultrahigh vacuum conditions) induces fast oxidation. The rate and depth of oxidation initially increase with increasing electron flux as indicated by results from Auger electron spectroscopy. Curves of oxygen development were fitted using a kinetic model that assumes the creation of oxide nucleation centers by the electron beam. The corresponding cross-sections were extracted. For fluxes exceeding 1 mA/cm2 , the oxidation rate is limited by the amount of oxygen present in the vacuum environment. For lower e-beam fluxes the oxidation process is slower with significant O chemisorption, resulting in shallower oxidation. These findings point out a way to control the thickness of nickel oxide in the nanometer range. Ó 2002 Elsevier Science B.V. All rights reserved. Keywords: Auger electron spectroscopy; Oxidation; Electron bombardment; Alloys; Nickel oxides
1. Introduction Oxidation phenomena of Ni3 Al alloys [1–4] are interesting from a technological viewpoint. The bulk material exhibits high thermal stability and it is protected from corrosion by an adherent Al2 O3 layer. Fundamental interest in Ni3 Al has arisen because it is a suitable model system for studying oxygen adsorption and the subsequent formation of oxides. Such studies are relevant for applications like exchange-bias (ferromagnetic/antiferromagnetic) junctions [5], heterogeneous metal/ oxide catalysts [6], aerospace technology [7,8], and lithography techniques in microelectronic device fabrication [9].
*
Corresponding author.
The nature of the oxide layers that are formed on Ni3 Al depends on the substrate temperature, as well as the partial pressures of oxygen and oxygencontaining elements. At room temperature (RT 300 K) the formation of Al2 O3 is hindered by the limited mobility of Al atoms in the highly ordered lattice. Instead oxygen is adsorbed at nickel sites or mixed Ni/Al sites. With increasing exposure O starts to be incorporated, leading to a disordered surface [1,2]. In a previous work [4] we demonstrated that the interaction with an electron beam strongly enhances nickel oxide growth on a Ni3 Al surface at RT and under ultrahigh vacuum (UHV) conditions. It is important to quantify this effect since the intense local oxidation interferes with surface processes in polycrystalline Ni3 Al (-B) alloys, such as segregation of boron [4]. For pure Ni, the
0039-6028/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 3 9 - 6 0 2 8 ( 0 2 ) 0 1 2 9 0 - 6
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
e-beam induced oxidation follows several stages: (i) dissociative chemisorption of oxygen, (ii) oxide nucleation and growth, and (iii) slow thickening of the oxide until saturation [10–15]. The second stage is particularly sensitive to the presence of nucleation sites around which oxide islands form and grow. These sites can be either intrinsic to the surface, or they can be created by incident electrons. Their metastable character ensures that short dosings with electrons are sufficient to induce rapid oxygen uptake [12]. The concept of nucleation centers for oxide growth is implicit in various analytical models that describe the stimulated oxidation of Ni [10–12,14]. For Ni3 Al(1 1 1) RT scanning tunneling microscopy (STM) studies also indicated the formation of small oxide islands on the surface [2]. Moreover, Ni is the primarily oxidizable constituent at this temperature [4]. Therefore in the present study we apply similar model descriptions to Ni3 Al. In fact, we focus on the influence of the electron flux density ue . The flux is varied both by changing the diameter d of the e-beam and the current I, since ue ¼ I=ðpd 2 =4Þ.
2. Experimental procedure Polycrystalline Ni3 Al samples were prepared by arc melting of the pure constituents with a small amount of boron (0.5–0.8 at.%). Boron is commonly added to reduce the inherent intergranular brittleness of this material. The overall composition was in the hypostoichiometric range with 24 at.% Al. The samples were homogenized at 1100 °C for 24 h to form the ordered L12 structure. Subsequently they were cross-sectioned, mechanically polished, and ultrasonically cleaned. In situ cleaning by Arþ sputtering occurred prior to Auger electron spectroscopy (AES) analysis. The AES system, with a base pressure of 4 108 Pa, is described in detail elsewhere [16]. The diameter of the e-beam can be varied in discrete steps of 5 lm, with a minimum diameter of about 15 nm at 10 keV accelerating voltage and 2.4 nA current. Auger mapping of oxygen was performed by recording the ratio (peak-background)/ background from the direct spectrum. The thick-
487
nesses of the local oxide layers were determined by depth profiling. Considering the fast rate of oxidation, it was not necessary to introduce additional oxygen into the system. As a result an exact quantification of the oxygen exposure in units of Langmuir is not possible. However, it is likely that molecular oxygen (O2 ) is not the only source of oxygen atoms. Another contribution might be due to the dissociation of molecules such as H2 O or CO. In general these are present in larger concentrations in a UHV atmosphere than O2 . In addition, chemical analysis of the samples was performed in a field ion microscope in atom probe mode (APFIM) [17]. Ions were removed from specially made Ni3 Al tips by field evaporation, and their flight times towards a detector were recorded. The time of flight for each ion is related to its mass-to-charge (m=n) ratio. Plots of ion number vs. their m=n ratio yielded detailed information on the material composition (Fig. 1).
3. Results and discussion Fig. 1 shows an APFIM time-of-flight spectrum of a Ni3 Al–B sample for which 5151 ions were collected in 211 min. The composition of the material was determined by ignoring ions from the ambient gas. The calculated values for Ni, Al, and B (Fig. 1, inset) are in good agreement with the relative amounts that were mixed when producing this material (input). Clearly O is absent from the spectrum, confirming our assumption that O only originated from the ambient atmosphere. Therefore the oxidation is not the result of O segregating from the bulk under the influence of the electron beam. Fig. 2 shows a series of e-beam stimulated oxidation measurements. The rate of oxidation is clearly dependent on the beam flux, which for the two spots is different by a factor of 25. An initial chemisorption stage cannot be identified. Similar behaviour has been observed for other materials with a high affinity for oxygen, where oxide nucleation begins before the saturation chemisorbed coverage is reached [18,19]. For Ni3 Al it is likely that both processes occur simultaneously, leading
488
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
Fig. 1. APFIM time-of-flight spectrum of a Ni3 Al specimen. The inset shows the measured elemental concentrations and the corresponding input values.
Fig. 2. Oxygen (508 eV) peak-to-peak Auger intensity vs. the oxidation time for e-beam current of I ¼ 2:4 nA. The inset shows a SEM image of exposed points (top) and Auger map of oxygen over the same area (bottom).
to a closely linear dependence as in the O-curve for d ¼ 50 lm. This curve shows a change in slope after 200 min marking the end of the chemisorption-dominated regime. We emphasize that the extent of oxidation in the surrounding areas was only a fraction of that in the exposed areas. This is evident from the
formation of bright spots as in the SEM image (inset of Fig. 2, top). An Auger map of O recorded over the same area demonstrates the local increase in oxygen intensity (inset of Fig. 2, bottom). The model proposed by Li et al. [11] that was used for fitting the O-curves can be expressed as hðtÞ ¼ hsat ðhsat hchem Þ exp½Kt ðk=ue rÞ
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
ðexpðue rtÞ 1Þ , where h is the oxide coverage, t is the exposure time, k is the rate constant for oxide growth, /e is the electron flux, r is the crosssection for the creation of nucleation centers, and hchem and hsat are the chemisorption saturation coverage and the oxide saturation coverage, respectively. The model assumes a completed chemisorption stage, i.e. h ¼ hchem at t ¼ 0. This assumption might be valid for the upper curve in Fig. 2 (d ¼ 10 lm) because any chemisorption is clearly negligible, but the same is not true for the lower curve (d ¼ 50 lm) since there is a significant contribution from chemisorption that extends beyond the initial stages. A modified model that includes a chemisorption term, reads in terms of Auger intensities: k I0 ðtÞ ¼ A ðA BÞ exp kt /e r fexp ð/e rtÞ 1g C expðkch tÞ ð1Þ with A the saturated Auger intensity, and B the chemisorption intensity. kch is now a rate constant for chemisorption. For the d ¼ 10 lm curve fit we set C ¼ 0. As the fits indicate the scenario of electroncreated nucleation centers for oxide growth appears to be valid for Ni3 Al. The faster rate for the largest beam flux translates into a larger value for the fitting parameter k, i.e. k ¼ 2:9 103 min1 , instead of 2:2 103 min1 for the lower flux. The chemisorption rate parameter kch for the latter turns out to be 6:1 103 min1 , which is almost three times larger than the corresponding value of k. The fits also yield the composite parameter /e r, which can be separated because the electron fluxes /e are known. The resulting cross-sections for creation of nucleation centers are r ¼ 3:7 1020 cm2 /electron (d ¼ 10 lm), and r ¼ 1:0 1020 cm2 /electron (d ¼ 50 lm). For pure Ni we previously obtained a much smaller cross-section of r ¼ 5 1022 cm2 /electron, for a beam with 5 lm diameter and remaining conditions the same [4]. The thicknesses of the local oxide layers were determined by performing a depth profile analysis (Fig. 3). A marked difference was found for the two electron flux conditions. The irradiation with
489
Fig. 3. Depth profile of oxygen acquired with a beam diameter of 5 lm.
0.1 mA/cm2 (d ¼ 50 lm) resulted in an oxide depth of 3:5 0:6 nm, compared to 6:0 0:6 nm for 3.0 mA/cm2 (d ¼ 10 lm). The uncertainty of 0.6 nm is equal to the escape depth of the 508 eV Auger electrons. Furthermore, oxidation measurements were performed with the e-beam current increased by an order of magnitude to I ¼ 24 nA (Fig. 4). The electron fluxes in this case were /e ¼ 30 and 1.2 mA/cm2 for the same beam diameters of d ¼ 10 and 50 lm, respectively. An additional spot of size d ¼ 70 lm (/e ¼ 0:6 mA/cm2 ) was included for comparison. For d 6 50 lm the oxidation occurs at a similar rate as before with the O-curves having a close resemblance. This indicates that the oxidation cannot proceed faster than a certain rate. At fluxes larger than 1 mA/cm2 the partial pressure of oxygen becomes insufficient to sustain the growth of nucleation sites. Estimations of the cross-sections from fits using Eq. (1) will then be inaccurate. At a given flux, it is expected that the lateral size of the beam also plays a role: the larger the exposed area, the more potential nucleation sites are covered, and the larger the contact area with the surrounding gas. Finally, in Fig. 4 the shape of the curve for d ¼ 70 lm is similar to that observed earlier for d ¼ 50 lm (Fig. 2). There is a relatively fast increase in oxygen intensity at the onset of exposure, followed by a region of more or less constant rate that eventually saturates. The
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S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
Fig. 4. Similar measurement as in Fig. 2 with e-beam current I ¼ 24 nA.
contribution of chemisorbed oxygen to the Auger signal is more significant than for the upper two curves.
Onderzoek der Materie’ (FOM) which is financially supported by the ‘Nederlandse Organisatie voor Wetenschappelijk Onderzoek (NWO)’.
4. Conclusions
References
We have demonstrated that the oxidation rate of Ni3 Al specimens and the oxide penetration depth strongly depend on the flux of incident electrons. For fluxes ue 6 0:5 mA/cm2 it is apparent that oxygen chemisorption occurs, while fast oxide growth dominates at larger fluxes. Increasing the flux beyond 1 mA/cm2 does not result in a faster oxidation rate, because the availability of oxygen in the UHV atmosphere becomes the limiting factor, rather than the number of created nucleation sites. At any rate, e-beam flux adjustment offers an alternative way to control nickel oxide thickness in the nanometer range.
[1] C. Becker, J. Kandler, H. Raaf, R. Linke, T. Pelster, M. Dr€ager, M. Tanemura, K. Wandelt, J. Vac. Sci. Technol. A 16 (1998) 1000. [2] A. Rosenhahn, J. Schneider, J. Kandler, C. Becker, K. Wandelt, Surf. Sci. 433–435 (1999) 705. [3] A. Rosenhahn, J. Schneider, J. Kandler, C. Becker, K. Wandelt, Appl. Surf. Sci. 142 (1999) 169. [4] S.A. Koch, D.T.L. van Agterveld, G. Palasantzas, J.Th.M. De Hosson, Surf. Sci. 476 (2001) L267. [5] J. Noguess, I.K. Shuller, J. Magn. Magn. Mater. 192 (1998) 203. [6] H.-J. Freund, H. Kuhlenbeck, V. Staemmler, Rep. Prog. Phys. 59 (1996) 283. [7] S.J. Sibener, R.J. Buss, C.Y. Ng, Y.T. Lee, Rev. Sci. Instrum. 51 (1980) 167. [8] P.N. Peters, H.C. Gregory, J.T. Swann, Appl. Opt. 25 (1986) 1290. [9] E.S. Snow, P.M. Campbell, F.K. Perkins, Naval Res. Rev. XLIX (1997) 15. [10] P.H. Holloway, J.B. Hudson, Surf. Sci. 43 (1974) 123, 141. [11] W. Li, M.J. Stirniman, S.J. Sibener, J. Vac. Sci. Technol A 13 (1995) 1574. [12] M.J. Stirniman, W. Li, S.J. Sibener, J. Chem. Phys. 103 (1995) 451.
Acknowledgements We would like to acknowledge financial support from the Netherlands Institute for Metals Research and the ‘Stichting voor Fundamenteel
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491 [13] W. Li, M.J. Stirniman, S.J. Sibener, Surf. Sci. 329 (1995) L593. [14] B.D. Zion, A.T. Hanbicki, S.J. Sibener, Surf. Sci. 417 (1998) L1154. [15] J.A. Slezak, B.D. Zion, S.J. Sibener, Surf. Sci. 442 (1999) L983. [16] D.T.L. van Agterveld, G. Palasantzas, J.Th.M. De Hosson, Acta Mater. 48 (2000) 195.
491
[17] W.A. Soer, master’s thesis, Materials Science Group, University of Groningen, 2001; W.A. Soer, P.M. Bronsveld, J.Th.M. De Hosson, Proceedings of IFES, Berlin, in press. [18] H. Brune, J. Wintterlin, J. Frost, G. Ertl, J. Wiechers, R.J. Behm, J. Chem. Phys. 99 (1993) 2128. [19] P.A. Thiry, J. Ghijsen, R. Sporken, J.J. Pireaux, R.L. Johnson, R. Caudano, Phys. Rev. B 39 (1989) 3620.
Electron beam induced oxidation of Ni3Al surfaces: electron flux effects S.A. Koch, G. Palasantzas, D.T.L. van Agterveld, J.Th.M. De Hosson
*
Department of Applied Physics, Materials Science Center and Netherlands Institute for Metals Research, University of Groningen, Nijenborgh 4, 9747 AG Groningen, Netherlands
Abstract Electron beam irradiation of polycrystalline boron doped Ni3 Al (at 300 K and under ultrahigh vacuum conditions) induces fast oxidation. The rate and depth of oxidation initially increase with increasing electron flux as indicated by results from Auger electron spectroscopy. Curves of oxygen development were fitted using a kinetic model that assumes the creation of oxide nucleation centers by the electron beam. The corresponding cross-sections were extracted. For fluxes exceeding 1 mA/cm2 , the oxidation rate is limited by the amount of oxygen present in the vacuum environment. For lower e-beam fluxes the oxidation process is slower with significant O chemisorption, resulting in shallower oxidation. These findings point out a way to control the thickness of nickel oxide in the nanometer range. Ó 2002 Elsevier Science B.V. All rights reserved. Keywords: Auger electron spectroscopy; Oxidation; Electron bombardment; Alloys; Nickel oxides
1. Introduction Oxidation phenomena of Ni3 Al alloys [1–4] are interesting from a technological viewpoint. The bulk material exhibits high thermal stability and it is protected from corrosion by an adherent Al2 O3 layer. Fundamental interest in Ni3 Al has arisen because it is a suitable model system for studying oxygen adsorption and the subsequent formation of oxides. Such studies are relevant for applications like exchange-bias (ferromagnetic/antiferromagnetic) junctions [5], heterogeneous metal/ oxide catalysts [6], aerospace technology [7,8], and lithography techniques in microelectronic device fabrication [9].
*
Corresponding author.
The nature of the oxide layers that are formed on Ni3 Al depends on the substrate temperature, as well as the partial pressures of oxygen and oxygencontaining elements. At room temperature (RT 300 K) the formation of Al2 O3 is hindered by the limited mobility of Al atoms in the highly ordered lattice. Instead oxygen is adsorbed at nickel sites or mixed Ni/Al sites. With increasing exposure O starts to be incorporated, leading to a disordered surface [1,2]. In a previous work [4] we demonstrated that the interaction with an electron beam strongly enhances nickel oxide growth on a Ni3 Al surface at RT and under ultrahigh vacuum (UHV) conditions. It is important to quantify this effect since the intense local oxidation interferes with surface processes in polycrystalline Ni3 Al (-B) alloys, such as segregation of boron [4]. For pure Ni, the
0039-6028/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 3 9 - 6 0 2 8 ( 0 2 ) 0 1 2 9 0 - 6
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
e-beam induced oxidation follows several stages: (i) dissociative chemisorption of oxygen, (ii) oxide nucleation and growth, and (iii) slow thickening of the oxide until saturation [10–15]. The second stage is particularly sensitive to the presence of nucleation sites around which oxide islands form and grow. These sites can be either intrinsic to the surface, or they can be created by incident electrons. Their metastable character ensures that short dosings with electrons are sufficient to induce rapid oxygen uptake [12]. The concept of nucleation centers for oxide growth is implicit in various analytical models that describe the stimulated oxidation of Ni [10–12,14]. For Ni3 Al(1 1 1) RT scanning tunneling microscopy (STM) studies also indicated the formation of small oxide islands on the surface [2]. Moreover, Ni is the primarily oxidizable constituent at this temperature [4]. Therefore in the present study we apply similar model descriptions to Ni3 Al. In fact, we focus on the influence of the electron flux density ue . The flux is varied both by changing the diameter d of the e-beam and the current I, since ue ¼ I=ðpd 2 =4Þ.
2. Experimental procedure Polycrystalline Ni3 Al samples were prepared by arc melting of the pure constituents with a small amount of boron (0.5–0.8 at.%). Boron is commonly added to reduce the inherent intergranular brittleness of this material. The overall composition was in the hypostoichiometric range with 24 at.% Al. The samples were homogenized at 1100 °C for 24 h to form the ordered L12 structure. Subsequently they were cross-sectioned, mechanically polished, and ultrasonically cleaned. In situ cleaning by Arþ sputtering occurred prior to Auger electron spectroscopy (AES) analysis. The AES system, with a base pressure of 4 108 Pa, is described in detail elsewhere [16]. The diameter of the e-beam can be varied in discrete steps of 5 lm, with a minimum diameter of about 15 nm at 10 keV accelerating voltage and 2.4 nA current. Auger mapping of oxygen was performed by recording the ratio (peak-background)/ background from the direct spectrum. The thick-
487
nesses of the local oxide layers were determined by depth profiling. Considering the fast rate of oxidation, it was not necessary to introduce additional oxygen into the system. As a result an exact quantification of the oxygen exposure in units of Langmuir is not possible. However, it is likely that molecular oxygen (O2 ) is not the only source of oxygen atoms. Another contribution might be due to the dissociation of molecules such as H2 O or CO. In general these are present in larger concentrations in a UHV atmosphere than O2 . In addition, chemical analysis of the samples was performed in a field ion microscope in atom probe mode (APFIM) [17]. Ions were removed from specially made Ni3 Al tips by field evaporation, and their flight times towards a detector were recorded. The time of flight for each ion is related to its mass-to-charge (m=n) ratio. Plots of ion number vs. their m=n ratio yielded detailed information on the material composition (Fig. 1).
3. Results and discussion Fig. 1 shows an APFIM time-of-flight spectrum of a Ni3 Al–B sample for which 5151 ions were collected in 211 min. The composition of the material was determined by ignoring ions from the ambient gas. The calculated values for Ni, Al, and B (Fig. 1, inset) are in good agreement with the relative amounts that were mixed when producing this material (input). Clearly O is absent from the spectrum, confirming our assumption that O only originated from the ambient atmosphere. Therefore the oxidation is not the result of O segregating from the bulk under the influence of the electron beam. Fig. 2 shows a series of e-beam stimulated oxidation measurements. The rate of oxidation is clearly dependent on the beam flux, which for the two spots is different by a factor of 25. An initial chemisorption stage cannot be identified. Similar behaviour has been observed for other materials with a high affinity for oxygen, where oxide nucleation begins before the saturation chemisorbed coverage is reached [18,19]. For Ni3 Al it is likely that both processes occur simultaneously, leading
488
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
Fig. 1. APFIM time-of-flight spectrum of a Ni3 Al specimen. The inset shows the measured elemental concentrations and the corresponding input values.
Fig. 2. Oxygen (508 eV) peak-to-peak Auger intensity vs. the oxidation time for e-beam current of I ¼ 2:4 nA. The inset shows a SEM image of exposed points (top) and Auger map of oxygen over the same area (bottom).
to a closely linear dependence as in the O-curve for d ¼ 50 lm. This curve shows a change in slope after 200 min marking the end of the chemisorption-dominated regime. We emphasize that the extent of oxidation in the surrounding areas was only a fraction of that in the exposed areas. This is evident from the
formation of bright spots as in the SEM image (inset of Fig. 2, top). An Auger map of O recorded over the same area demonstrates the local increase in oxygen intensity (inset of Fig. 2, bottom). The model proposed by Li et al. [11] that was used for fitting the O-curves can be expressed as hðtÞ ¼ hsat ðhsat hchem Þ exp½Kt ðk=ue rÞ
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
ðexpðue rtÞ 1Þ , where h is the oxide coverage, t is the exposure time, k is the rate constant for oxide growth, /e is the electron flux, r is the crosssection for the creation of nucleation centers, and hchem and hsat are the chemisorption saturation coverage and the oxide saturation coverage, respectively. The model assumes a completed chemisorption stage, i.e. h ¼ hchem at t ¼ 0. This assumption might be valid for the upper curve in Fig. 2 (d ¼ 10 lm) because any chemisorption is clearly negligible, but the same is not true for the lower curve (d ¼ 50 lm) since there is a significant contribution from chemisorption that extends beyond the initial stages. A modified model that includes a chemisorption term, reads in terms of Auger intensities: k I0 ðtÞ ¼ A ðA BÞ exp kt /e r fexp ð/e rtÞ 1g C expðkch tÞ ð1Þ with A the saturated Auger intensity, and B the chemisorption intensity. kch is now a rate constant for chemisorption. For the d ¼ 10 lm curve fit we set C ¼ 0. As the fits indicate the scenario of electroncreated nucleation centers for oxide growth appears to be valid for Ni3 Al. The faster rate for the largest beam flux translates into a larger value for the fitting parameter k, i.e. k ¼ 2:9 103 min1 , instead of 2:2 103 min1 for the lower flux. The chemisorption rate parameter kch for the latter turns out to be 6:1 103 min1 , which is almost three times larger than the corresponding value of k. The fits also yield the composite parameter /e r, which can be separated because the electron fluxes /e are known. The resulting cross-sections for creation of nucleation centers are r ¼ 3:7 1020 cm2 /electron (d ¼ 10 lm), and r ¼ 1:0 1020 cm2 /electron (d ¼ 50 lm). For pure Ni we previously obtained a much smaller cross-section of r ¼ 5 1022 cm2 /electron, for a beam with 5 lm diameter and remaining conditions the same [4]. The thicknesses of the local oxide layers were determined by performing a depth profile analysis (Fig. 3). A marked difference was found for the two electron flux conditions. The irradiation with
489
Fig. 3. Depth profile of oxygen acquired with a beam diameter of 5 lm.
0.1 mA/cm2 (d ¼ 50 lm) resulted in an oxide depth of 3:5 0:6 nm, compared to 6:0 0:6 nm for 3.0 mA/cm2 (d ¼ 10 lm). The uncertainty of 0.6 nm is equal to the escape depth of the 508 eV Auger electrons. Furthermore, oxidation measurements were performed with the e-beam current increased by an order of magnitude to I ¼ 24 nA (Fig. 4). The electron fluxes in this case were /e ¼ 30 and 1.2 mA/cm2 for the same beam diameters of d ¼ 10 and 50 lm, respectively. An additional spot of size d ¼ 70 lm (/e ¼ 0:6 mA/cm2 ) was included for comparison. For d 6 50 lm the oxidation occurs at a similar rate as before with the O-curves having a close resemblance. This indicates that the oxidation cannot proceed faster than a certain rate. At fluxes larger than 1 mA/cm2 the partial pressure of oxygen becomes insufficient to sustain the growth of nucleation sites. Estimations of the cross-sections from fits using Eq. (1) will then be inaccurate. At a given flux, it is expected that the lateral size of the beam also plays a role: the larger the exposed area, the more potential nucleation sites are covered, and the larger the contact area with the surrounding gas. Finally, in Fig. 4 the shape of the curve for d ¼ 70 lm is similar to that observed earlier for d ¼ 50 lm (Fig. 2). There is a relatively fast increase in oxygen intensity at the onset of exposure, followed by a region of more or less constant rate that eventually saturates. The
490
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491
Fig. 4. Similar measurement as in Fig. 2 with e-beam current I ¼ 24 nA.
contribution of chemisorbed oxygen to the Auger signal is more significant than for the upper two curves.
Onderzoek der Materie’ (FOM) which is financially supported by the ‘Nederlandse Organisatie voor Wetenschappelijk Onderzoek (NWO)’.
4. Conclusions
References
We have demonstrated that the oxidation rate of Ni3 Al specimens and the oxide penetration depth strongly depend on the flux of incident electrons. For fluxes ue 6 0:5 mA/cm2 it is apparent that oxygen chemisorption occurs, while fast oxide growth dominates at larger fluxes. Increasing the flux beyond 1 mA/cm2 does not result in a faster oxidation rate, because the availability of oxygen in the UHV atmosphere becomes the limiting factor, rather than the number of created nucleation sites. At any rate, e-beam flux adjustment offers an alternative way to control nickel oxide thickness in the nanometer range.
[1] C. Becker, J. Kandler, H. Raaf, R. Linke, T. Pelster, M. Dr€ager, M. Tanemura, K. Wandelt, J. Vac. Sci. Technol. A 16 (1998) 1000. [2] A. Rosenhahn, J. Schneider, J. Kandler, C. Becker, K. Wandelt, Surf. Sci. 433–435 (1999) 705. [3] A. Rosenhahn, J. Schneider, J. Kandler, C. Becker, K. Wandelt, Appl. Surf. Sci. 142 (1999) 169. [4] S.A. Koch, D.T.L. van Agterveld, G. Palasantzas, J.Th.M. De Hosson, Surf. Sci. 476 (2001) L267. [5] J. Noguess, I.K. Shuller, J. Magn. Magn. Mater. 192 (1998) 203. [6] H.-J. Freund, H. Kuhlenbeck, V. Staemmler, Rep. Prog. Phys. 59 (1996) 283. [7] S.J. Sibener, R.J. Buss, C.Y. Ng, Y.T. Lee, Rev. Sci. Instrum. 51 (1980) 167. [8] P.N. Peters, H.C. Gregory, J.T. Swann, Appl. Opt. 25 (1986) 1290. [9] E.S. Snow, P.M. Campbell, F.K. Perkins, Naval Res. Rev. XLIX (1997) 15. [10] P.H. Holloway, J.B. Hudson, Surf. Sci. 43 (1974) 123, 141. [11] W. Li, M.J. Stirniman, S.J. Sibener, J. Vac. Sci. Technol A 13 (1995) 1574. [12] M.J. Stirniman, W. Li, S.J. Sibener, J. Chem. Phys. 103 (1995) 451.
Acknowledgements We would like to acknowledge financial support from the Netherlands Institute for Metals Research and the ‘Stichting voor Fundamenteel
S.A. Koch et al. / Surface Science 507–510 (2002) 486–491 [13] W. Li, M.J. Stirniman, S.J. Sibener, Surf. Sci. 329 (1995) L593. [14] B.D. Zion, A.T. Hanbicki, S.J. Sibener, Surf. Sci. 417 (1998) L1154. [15] J.A. Slezak, B.D. Zion, S.J. Sibener, Surf. Sci. 442 (1999) L983. [16] D.T.L. van Agterveld, G. Palasantzas, J.Th.M. De Hosson, Acta Mater. 48 (2000) 195.
491
[17] W.A. Soer, master’s thesis, Materials Science Group, University of Groningen, 2001; W.A. Soer, P.M. Bronsveld, J.Th.M. De Hosson, Proceedings of IFES, Berlin, in press. [18] H. Brune, J. Wintterlin, J. Frost, G. Ertl, J. Wiechers, R.J. Behm, J. Chem. Phys. 99 (1993) 2128. [19] P.A. Thiry, J. Ghijsen, R. Sporken, J.J. Pireaux, R.L. Johnson, R. Caudano, Phys. Rev. B 39 (1989) 3620.