Light impurities in oxidizing metal foils

Materials Science and Engineering, A125 (1990) 235-240

235

Light Impurities in Oxidizing Metal Foils D. FINK and K. TJAN* Hahn-Meitner-lnstitut, Glienickerstrasse100, D-tO00Berlin 39 (F.R.G.) M. F. DA SILVA LaboratorioNacionalde Engenhariae TechnologiaIndustrial, P-2686Sacavdm (Portugal) J. C. SOAREZ Centro de FisicaNuclear, Universidadede Lisboa, P-1699Lisboa Codex (Portugal) (ReceivedMay 1, 1989; in revisedform October2, 1989)

Abstract Distributions of implanted lithium ions and of natural boron impurities are monitored in titanium foils after isochronal annealing in high vacuum. A strong correlation between the shape of the depth profiles and the oxide growth in titanium is observed. In contrast, the depth distribution of lithium implanted into zirconium remains nearly unaltered up to high temperatures. The observed oxide growth of titanium in air at elevated temperatures is twice as fast as reported earlier. The activation energy for the oxide growth, 1.2 eV, could, however, be confirmed. Annealing in an oxygen-deficient atmosphere not only results in a retarded oxide growth rate but also exhibits a different activation energy (0. 7 eV), which indicates different chemical oxidation reactions. 1. Introduction

The diffusive behaviour of fight impurities, such as lithium or boron, in homogeneous matter during thermal annealing has been frequently studied, because of its great technological consequences. However, impurity redistribution in the presence of a moving boundary, e.g. a growing oxide layer, has only rarely been examined. Such work has been devoted nearly exclusively to the Si/SiO 2 system, for electronic applications. Other systems, e.g. metal/oxide structures, such as Ti/TiO 2 or Zr/ZrO2, containing light impurities such as lithium or boron, have not yet been examined. This is the subject of this work. *Present address: I1. Tambakan 32, Muntilan 56412, Jateng, Indonesia. 0921-5093/90/$3.50

We chose titanium and zirconium for their technological importance. Furthermore, a comparison of the behaviour of both materials is interesting, in so far as the metal and oxide phases for zirconium and titanium (h.c.p. and tetragonal respectively) are similar for both systems. Both metals exhibit extensive oxygen solid solubility and a compact nature of their oxides. Some of the results presented here were obtained earlier [1 ]. 2. On the thermal oxidation of undoped titanium and zirconium

As there has already been some work [2-4] performed on the thermal oxidation of pure titanium and zirconium foils, we first present an overview about the present knowledge. Oxygen depth profiles in Ti/TiO2 or Zr/ZrO2 as well as Si/titanium silicide/TiO 2 systems have been studied by thermal annealing in normal air and characterized by thermogravimetry, microscopy and microhardness measurements [2], or by electron probe microanalysis and Rutherford backscattering spectroscopy (RBS) [3, 4]. From these results, it follows that the oxidation process leads to (1) a top oxide layer with relatively welldefined stoichiometry and (2) a subsequent underlying oxygen-enriched zone, with a long exponentially shaped concentration profile, until (3) the natural oxygen impurity concentration in the bulk is reached. No oxide precipitation has been observed to occur in the bulk. The thickness z (cm) of the oxide layer is given in the case of titanium [2] by the relation { 27"350 (cal m°l-~) } Zox= 1.5t exp RT

(1)

© ElsevierSequoia/Printedin The Netherlands

236 with R (=1.987 cal mo1-1 K -1) is the gas constant, t (s) is the annealing time and T (K) is the temperature. The initial concentration Cox of the exponential tail at the oxide-metal interface is determined by the temperature-dependent oxygen diffusion constant and the diffusion time, typical values for titanium being 14-34 at.% in the temperature range 650-900 °C after annealing [2] for 50-100 h. The width of the oxygen tail (measured from the oxide-metal interface towards the bulk) can be well described by the depth Z5at.% at which the oxygen concentration reaches the value 5 at.%. (This value shows up in the microscopic studies [2] by a step in reflectivity of the etched titanium surface, for an unknown reason.) Typical values of Z5at.% reported in ref. 2 for titanium range from about 5 g m at 600°C to about 45 g m at 750 °C, after annealing times of the order of 100 h. For zirconium [4], the corresponding values are roughly comparable with those for titanium. It should be pointed out that all these results have been obtained only for air under normal pressure. No results have yet been obtained for oxidation in oxygen-deficient atmospheres.

3. Experimental procedure As the metals under discussion easily absorb impurities which might alter their properties, we decided first to check the above-mentioned results by some of our own measurements. We restricted ourselves to the more reactive metal, titanium. As the nominal purity of the titanium foils as supplied by the producer (MRC) is 99.97 at.%, we hence have to assume a dissolved oxygen component of around some 0.01 at.%. The titanium foils were cleaned with acetone in an ultrasonic bath but not treated otherwise. Then, in a first experimental series, the titanium foils were annealed isochronally in steps of 30 min, in air at atmospheric pressure, and analysed by RBS with a 1.5 MeV proton beam at the 2 MeV Van de Graaff accelerator of the Laboratorio Nacional de Engenharia e Technologia Industrial, Sacav6m. The layer thicknesses were determined with an accuracy of _+20%. Apart from pure titanium foils, lithiumimplanted titanium and zirconium foils were also examined. They were prepared by 6Li ÷ implantation at the Danfysik heavy ion accelerator SIB of the Hahn-Meitner-Institut, Berlin, at energies

between 50 and 300 keV and current densities of typically 0.1 # A c m -2 for doses between 5 x 1014 and 30 x 1014 ions cm -2, after having undergone the same cleaning treatment as above. The total pressure during implantation was ( 2 - 3 ) x 10 -6 Torr. Annealing of the implanted samples was performed for 1 h, in steps of 50 or 100 °C under dry argon gas at a pressure of about 2 x 10-6 Torr (the oxygen partial pressure was estimated by means of a residual gas analyser to be of the order of 10 -8 Torr). During annealing, the total pressure rose to ( 1 - 5 ) x 10 -4 Torr, essentially owing to degassing of the surrounding molybdenum walls of the furnace. The lithium depth profile measurements were performed at our neutron depth profiling facility, S-30, at the high flux reactor of the Institute Max von Laue-Paul Langevin, Grenoble, France, where thermal neutrons were used as projectiles to apply the nuclear reaction 6Li(n, ct)aH for depth profiling. This method has been described extensively earlier (see for example ref. 5). In brief, the energy loss, which the monoenergeticaUy created a particles suffer during their passage from the depths of their origins towards the sample's surface, is measured by use of a semiconductor detector, then analysed in a multichannel analyser and finally converted to the original depth information by a subsequent computer program. This program makes use of the relation

6x =

diE cos ct °

s(e)

(2)

6E being the a particle's energy loss during the passage of the depth interval 6x, the particle's direction being tilted by a ° from the sample's normal, and S(E) being the a particle stopping power in the sample medium. S(E) is known within _+5% for most target materials [6], and the depth was determined t o within _ 213 ,/k for titanium and to within _+191 ,/k resolution for zirconium, using a surface barrier detector with 17 keV full width at half-maximum (for 2 MeV a particles). The sensitivity is typically 10 -6 (1 at.ppm). 4. Results and discussion Our RBS results on the oxide growth of pure titanium foils at atmospheric pressure (full din-

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monds in Fig. 1 ) are higher than those reported in the literature ([2], open circles in Fig. 1) by a factor of about 2 but otherwise exhibit a similar behaviour. This is shown in Fig. 1, where the oxide growth is illustrated by the generalized plot of log( Zox/ t 1/2) vs. T - 1, as was done in ref. 2 (i.e. a modified Arrhenius plot). From this, we deduce agreement, in principle, with the titanium properties reported earlier. Figure 2 shows the measured lithium depth distributions in titanium (full circles) for different annealing stages, and Fig. 3 gives the corresponding results for lithium in zirconium. As the thermal neutrons exhibit a high reaction cross-section not only for 6Li but also for l°B, we can measure these nuclides simultaneously, when they are present in sufficient concentrations. In fact we found quite a remarkable amount of boron emerging during subsequent annealing, the depth distribution of which undergoing various stages at increasing temperatures, similar to the implanted lithium. These boron distributions are included in the figures (as open circles). As there is no external source for boron during the annealing procedure, we have to ascribe the origin of these boron depth profiles to boron impurities in the material. Apparently, boron becomes mobile at elevated temperatures and is preferentially incorporated in the oxide layer. The accumulated maximum concentrations are of the order of 0.1-1 at.% and cannot be presumed to lead to the formation of a new phase in the absence of evidence to that effect. Surface enrichment of boron impurities is a very common feature and has been found previously for many other metals--implanted as well as unimplanted--

e.g. for nickel, platinum, niobium and silver (see also ref. 8). Close observation of both the lithium and the boron depth profiles in titanium (Fig. 2) reveals that all distributions exhibit some irregularity, e.g. a change in slope, or even a peak, at a specific temperature-dependent depth. We tend to believe that these irregularities reflect an oxygen-metal interface. In fact, the residual gas at which these implanted metal foils were annealed (2 x 10-65 × 1 0 - 4 Torr total pressure; see above) has still to be regarded as a somewhat oxidizing atmosphere. In support of this, the work of Merchant and Amano [3] is cited, in which it was found that titanium evaporation at 2 × 1 0 - 7 Torr already resulted in an oxygen uptake of up to 30 at.%. Similarly, using a different measuring technique, Unnam et al. [2] found that the total air pressure has to be lower than 7 × 1 0 - 7 Torr to prevent oxygen uptake of titanium foils. The observed irregularities in the depth profiles at the metal-oxide interface positions are quite strong for low annealing temperatures and somewhat smoothed out for higher annealing temperatures. We explain this by the increasing concentration of dissolved oxygen beyond the oxide layer at elevated temperatures, which renders the physical environment for the impurities in the oxide and the metallic matrix less different. If we deduce the probable interface positions Zox from Fig. 2 and plot them on the modified Arrhenius plot of log(Zox/t 1/2) vs. T -l, as was done earlier for annealing at atmospheric pressure (Fig. 1, full circles), we get another straight line, but lower and with a smaller slope than the previous line. The difference in heights has to be ascribed to the different abundances of oxygen in the experiments. The difference in the slopes, which reflects the different activation enthalpies (1.2 eV for oxidation in normal air, and 0.7 eV for oxidation under high vacuum conditions), indicates that different chemical reactions occurred in the two cases. It might be speculated that in an oxygen-deficient atmosphere, TiO or Ti203 is formed, whereas at normal air the more stable TiO 2 is created. In detail, the evolution of the lithium and boron depth distributions in titanium might be explained as follows. (a) From room temperature to 450 °C, regular lithium implantation profiles are demonstrated, which can be well described by the present theory

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Fig. 2. Depth profiles of implanted lithium (e) and of precipitated boron (o) in titanium. The oxygen-metal interface, as deduced from the impurity profiles, is indicated by arrows. (a) A comparison with the theoretically expected range profile (histogram according to the T R I M code, see refs. 1 and 7) and its convolution with the detector resolution (smooth line) are given additionally. (b) The corresponding theoretical nuclear damage distribution (smooth line, not convoluted). Depth profiles are shown for various temperatures T: curve a, 2 0 - 4 5 0 °C; curve b, 5 0 0 - 6 0 0 °C; curve c, 700 °C; curve d, 800 °C; curve e, 900 °C; curve f, 1000 °C.

[7, 9] (see the histogram in Fig. 2(a)). No surface segregation is found, indicating zero lithium mobility after implantation. The boron content in the sample cannot be distinguished from the background value in the measuring chamber, and hence is assumed to be of the order of 1 at.ppm or less--except for a slight near-surface enrichment of about 10 at.ppm which was also present in the unimplanted samples. (b) After annealing at 500-600°C, lithium exhibits some mobility and tends to redistribute from the range to the nuclear damage distribution

(see the smooth curve in Fig. 2(b), which also indicates a slight shift of the profile maximum towards the surface). Simultaneously, some lithium accumulates at the surface, prol°~. 'y owing to bonding with oxygen in the evolving oxide layer. Apparently, boron also becomes mobile at this temperature and is preferentially trapped at the surface or incorporated in the underlying oxide layer. (c) At 700°C, the lithium surface peak has vanished by sublimation. The depth profile in the bulk has conformed to the shape of the nuclear

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damage profile nearly perfectly (i.e. the radiationinduced defects are still stable). The boron surface concentration still increases. (d) At 800 °C, the lithium depth profile shows a dramatic depletion in the oxide region, although the bulk distribution remains unaltered. Most of

the surface boron is transported into the bulk along the moving oxide-metal interface, and only a minor fraction remains in the oxide. The solubility ratio of oxide and metal is found to be less than 10-2 for lithium, and about 0.4 for boron. (e) At 900 °C, the oxide-metal interface passes

240

the region of maximum concentration of the bulk lithium depth distribution, which is moved somewhat towards the interior. Simultaneously, lithium starts to diffuse deeply into the metal matrix. Boron behaves similarly. (f) At 1000 °C, finally, we find a smoothing out of the lithium and boron structures, the profiles' maxima now being entirely submerged in the oxide layer. This means that the oxide layer growth is faster that the lithium and boron diffusion, i.e. the moving oxide-metal interface front "overtook" the moving impurities. Interestingly, the lithium tail extending into the bulk did not broaden, which indicates that lithium no longer diffuses strongly; rather it is bonded to traps, presumably forming Li20 with the highly abundant oxygen. All our experiments with lithiumimplanted titanium foils yielded essentially the same results, independently of the implantation fluence ((5-30)x 1014 ions cm -2 s-1) and energy (50-300 keV) and hence the depth (0.1-1/~m). In spite of exhibiting some similarity with lithium in titanium at lower temperatures, lithium in zirconium behaves in a strikingly different way at higher temperatures, see Fig. 3. After implantation, we find a regular lithium range profile, without any surface enrichment. After annealing up to 900 °C, the profile shape remains essentially unchanged, except for some broadening on the far side of the surface. This may indicate the onset of lithium mobility. Some lithium also accumulates in the oxide layer, but not at the surface itself (where it would sublime immediately at this high temperature). In contrast with titanium, the moving oxide-metal interface does not affect the impurity profiles strongly enough to enable us to identify it.

5. Conclusions The oxide growth of titanium foils at elevated temperatures in air was found to proceed twice as fast as reported earlier by Unnam et al. [2]. A deficiency in oxygen abundance results not only in a decrease in the oxide growth velocity

but also in a different activation enthalpy. This indicates that different chemical reactions are occurring, presumably the formation of TiO and Ti203 instead of TiO 2. Depth distributions of implanted lithium and of natural boron impurities in titanium are markedly affected by the growing oxide layer, so that the position of the oxide-metal interface can be derived from the lithium and boron depth profiles. This is, however, not possible for lithium and boron in zirconium, in spite of some similarities in the lithium and boron depth distributions in titanium and zirconium at low temperatures.

Acknowledgments This work was made possible by the kind support of the Institute Laue-Langevin, Grenoble, France, the Deutscher Akademischer Austauschdienst and the Internationales BiJro of the Kernforschungsanlage Karlsruhe, ER.G. We thank Mr. M. Bri6re for valuable discussions.

References 1 Tjan Kie, Thesis, Free University, Berlin, 1985. 2 J. Unnam, R. N. Shenoy and R. K. Clark, Oxid. Met., 26 (3-4) (1986) 231, and refs. 1-11, 14 and 15 cited therein. 3 P. Merchant and J. Amano, J. Vac. Sci. Technol. B, 2 (4) (1984) 762. 4 M. Berti, A. Camera, A. V. Drigo, A. Armigliato, A. Desalvo and R. Rosa, Nucl. lnstrum. Methods, 182-183 (1981)215. 5 J. P. Biersack, D. Fink, R. Henkelmann and K. Miiller, Nucl. lnstrum. Methods, 149 (1978) 93. 6 J. E Ziegler, J. P. Biersack and U. Littmark, Proc. U.S.-Japan Semin. on Charged Particle Penetration Phenomena, in ORNL Rep. CONF-820131, 1982, p. 88 (Oak Ridge National Laboratory). 7 J. P. Biersack and L. G. Haggmark, Nucl. lnstrum. Methods, 174 (1980) 257. 8 D. Fink, Mater. Sci. Eng., Al15 (1989) 89. 9 D. Fink, Tjan Kie, J. P. Biersack, Wang Lihong and Ma Yunru, submitted to Radiation Effects and Defects in Solids, 1989.