External oxidation of aluminium-lithium alloys

Surface and ComingsTechm)iogy. 51 (1992) 57-64

57

External oxidation of aluminium-lithium alloys Dietmar Fink Hahn-Meitner-institut. Dell. P-3. Gliee6ckerstrasse I00. W-IO00 Berlin 39 ( FRG )

V. Hnatowicz, J. Kvitek and V, Havranek Ceskoslot~easko Akademia Vod. Ustae J¢~derneFy2vky. 25068 ~e~ near Prague ( Czechoslor~akia)

J. T. Zhou* Hatm-Meitner-lnstitut, Dept, P-3. Giienickerstrasse I00. W-IO00 Berlin 39 (FAG)

Akstraet The thermal oxidation dynamics of cold rolled aluminium-3%lithium foils was studied as a function of the annealing temperature, time and environmental gas pressure, by means of Rutherford backscattering spectrometry, thermal neutron depth profiling and scanning electron microscopy, Highly complex behaviour was found, which includes the depletion of lithium atoms from sub-surface layers and their enrichment in the surface area as oxide, The ratio oflithium to oxygen is usually far from stoichiometric. The surface transformation into oxide proceeds very rapidly in the first few minutes, and subsequently hardly shows any change,

1. latrmktetiea Hitherto, relatively little knowledge has been gained about the oxidation kinetics of Li-A1 alloys. We studied this process in some more detail. Lithium-aluminium alloys have gained considerable comraercial interest over several decades as lightweight high strength construction materials [11, for example in the aerospace industry [2-51, as anode materials in high-performance electrolyric cells such as LiAI/FeS~ (1 ~
*On leave from Shanghai Institute of Metallurgy, Ion Beam Laboratory, Academia Sinica, Shanghai 200050. China,

0257-8972/92/$5,00

Our techniques of examination cover Rutherford backscattering spectrometry (RBS), thermal neutron depth profiling (NDP) and scanning electron microscopy (SEM). With RBS, we essentially study the behaviour of the heavier matrix components aiuminium and oxygen, the NDP technique is specific for the lithium component of the alloy and for eventual boron contamination, and the surface appearance during oxidation is monitored by SEM. The results of this work are compared with the earlier findings of titanium and zirconium oxidation.

2. Exper~memai Iletaiis The A1-Li alloy was prepared by the following procedure. To prevent oxidation, the lithium (90% enriched in lithium-6) was handled under paraffin oil. The paraffin was then replaced by n-pentane, thus combining a protective function and high volatility. The lithium sample was transferred inside a crucible of the host aluminium material and inserted into a high vacuum induction furnace. After evaporation of the protective npentane, the melting process was performed under an atmosphere of purified argon [221, The prepared alloy was then analysed chemically and with NDP for its lithium content, both techniques yielding about 3 at.% lithium. Subsequently, foils were produced by cold roiling the material down to thicknesses of 100 and 10 l~m. These foils were earlier used to study 3 H and

(~ 1992 - - Elsevier Sequoia, All rights reserved

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O.~'idalio+~OI AI Li alloy,~

stopping +R+wcrs in ki AI alloys [22], and to explore lithium clustering by positron annihilation [23]. In spite of storage of the foils in normal air at r¢R+mtemperature for about 6 years, their visual appearance was still gotRt. Before using the alloys again, their surface was etched with HCI, and cleaned with water and acetone+ Then, the samples were subjected to a systematic annealing treatment with variation of parameters: annealing temperature 150, 100, 150..... up to 550 C), time 18, 16, 32 s i n , I, 2, 4, 8 h ..... up to about 16(X)hi, and ambient air pressure (760, 10 "- and 10 ~ Torr, controlled by ;i manometer). For each treatment, a new piece of foil was used, all foils originating from the same batch. The samples were then analysed to study the changes in depth profiles of the alloy co.mlxments aluminium and lithium, and of in-diffUsed oxygen from the air. These examinations were performed at the l~e~.institute by RBS using a 2,35 MeV proton beam impinging perpendicularly into the samples to study their overall appearance, and a 2.0 MeV '~He ion beam in the same direction to probe details of the samples' surfaces. The detector was placed at an angle of 160 to the analysing beam. For calibration of the relative proton elastic scattering cross-sections, we used an Li,CO.~ standard. At 2.35 MeV, we derived a cross-section ratio a(Li):o(O):o(C)=l :1.1+0.1:1.2+0.1. The lithium cross-section is a slowly decreasing function of the proton energy but the oxygen and carbon cross-sections are nearly constant in the relevant energy region. Additionally, NDP studies were performed using the high flux reactor (HFR) of the Institute Laue-Langevin fiLL), Grenoble. This technique uses the thermal neutron-induced nuclear reaction 6Li(n,0t)-~H to study the depth distribution of lithium, it has the advantage of high sensitivity and large profiling depth, and is also a non-destructive technique. The basic principles of this technique, as well as physical and technical details, have been published previously [24-26]. Finally, we took a number of (SEMI images to enable a reliable interpretation of the depth protile measurements.

3, Thermd mmeadlimgof A I - L i alloys: remd(+sa,ml

3.1. As-prepared samples Measurements performed by NDP immediately after the preparation procedure in 1984 revealed a homogeneous lithium distribution up to the sample's surface (Fig. I, curve for 20 +C). Repeating the same analysis 6 years later did not show any change in the lithium profile, i.e. we have to interpret this as negligible lithium mobility in the alloy at room temperature. Additional

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RBS examinations al,,a) showed that the aluminium distribution is homogeneous, and that no significant superficial or bulk oxygen has accumulated (Fig. 2. curves for oxygen). In contrast. Auger electron spectroscopy (AES) examinations always show prom~uncexJ oxygen peaks for as-prepared aluminium and commercial AI Li alloy films [20].

3.2. Annealing in high t~acuum Interestingly, isochronal annealing up to 350 C for I h each did not result in any depth profile changes, independent of the environmental conditions (dry air. prevacuum o f a ~ u t 5 × l0 z Tort, high vacuum of about l0 ~ Torr). This is remarkable insofar as the annealing of lithium-implanted aluminium samples (at high vacuum. l0 ~ Torr) showed depth profile changes at temperatures as low as 250 C, and strong lithium mobility after 350 C annealing for only 30 s i n [27. 28]. This difference in mobility may be the result of a combined action of the presence of mobile radiation-induced wlcancics and interstitiais in the implanted material on the one hand, and on the other hand the higher abundance of oxygen in the surface layer of the unimplanted alloy, which tends to bond the lithium chemically, thus reducing its near-surface mobility. Whereas Jardin and Robert [20] reported the onset of lithium mobility in commercial AI- Li- Mg alloys only at temperatures above 430 C in ultrahigh vacuum (around 10 -~ Torr), we found the first changes in the depth profiles already above 350 C . The oxidation

D, Fiaket at, / Oxidmion of AI-Li alloys

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spectra of an unanncaled sample (curve 0), after 2 h (curve I), and after 3.5 h anneal at 500 ~'C in prcvacuum (10 -z Tort) (curve 2),

proceeded much slower under high vacuum conditions than at higher pressures (s¢¢ below), probably owing to the reduced oxygen supply. The results suggest that first lithium becomes mobile in the alloy and eventually reaches the surface. This is confirmed by Jardin and Robert [20]. According to their observations, the migration of lithium towards the surface preferentially occurs through the AlzOs grain boundaries, which act as short circuits for the diffusion. At those relatively low temperatures, lithium diffusion by either substitutional or interstitial mechanisms through the AlzO3 structure is still impossible. It appears that the migration takes place in the form of lithium oxide rather than as metallic lithium

(this state is not detected by electron loss spectroscopy (ELS) [20]), the oxygen stemming (according to rd. 20) from a reduction of the A1203 at the Al203-A! interface. ELS and AES [20] show that aluminium is not present in the Li20 surface precipitation zone. The superficial enrichment of lithium implies lithium depletion in underlying layers. This is actually found. In the initial annealing stages, the lithium concentration profiles in the depiction zone are well described by the theoretically expected error-function shape, see Fig. I. At a later stage, this zone becomes broader, with some constant minimum lithium concentration (see also ref. 18).

D, Fink et al,

60

Oxidation ¢)I 41 l.i alloys

3.3. Annealing in preracuum

With increasing oxygen abundance, the oxide growth increases dramatically. Figure 2 illustrates the details of growth of the ,superficial oxide layer, aceornpanied by a superficial depletion in aluminium and an enrichment in lithium. In the initial stage of annealing, both aluminium and lithium are present in the near-surface region (see, for example, Fig. 2, curves I), With increasing lithium atoms arriving, the aluminium is gradually replaced, to produce a lithium oxide layer on the surface (,see, for example, Fig. 2, curves 2). Table I presents an overview of the parameters of the oxygen distributions derived from the RBS spectra for isothermal annealing at 500 C . Figure 3(a) shows typical examples of the oxygen distributions which we found. They exhibit a peak at the surface, and a more or less exponential decrease in concentration down to larger depths. We have characterized them by their width, .see Table I. It appears that this superficial oxide structure builds up very rapidly, within the first few minutes, and then is more or less stationary for the next few days, with little correlation between the annealing conditions and the measured depth profile shapes. The typical profile width is around 100 nm. Only after prolonged times, of the order of months, are some slight changes in the oxygen depth profile "shapes visible. As an example, ,see Fig. 3(b), which shows a three times broader oxygen profile than Fig. 3(a). Owing to our present measuring geometry, we can exclude the possibility that the observed oxygen distribution is due to surface roughness. We cannot exclude, however, the formation of a nonuniform oxide layer. From the oxygen structures, the thickness of the

lithium oxide layers can be estimated to be 0.12 ~tm alter 2 h, and 0.25 lam after 3.5 h annealing time. The total oxygen areal density was determined from the measured proton spectra relative to the above-mentioned SiO, standard (see Fig. 3), comprising 5,5 × 10~ atoms cm -~. The areal oxygen densities do not show any significant correlation with the annealing time, once the oxide layer has been established, see Table I. Only on a few of the 5 0 0 C annealed samples could the lithium signal be verified in the RBS s ~ t r a . In the other cases, we used information derived from N D P measurements. The corresponding atomic [Li],:[O] ratios are also compiled in Table I. In all cases (except one), the ratio is far from Li,O stoichiometry, it ,seems that, independent of the annealing conditions, the [Li]. [O] ratio is relatively constant around 0.3+0.3, the individual fluctuations being astonishingly large. This implies that the excess oxygen in our samples is bonded to ,some other element (perhaps to aluminium or to carbon), or not bonded at all. Only for prolonged annealing times of the order of months is the correct stoichiometry partly reached, see Fig. 3(b). The measurements with 2.0 MeV ~ particles clearly show that the surface structure, as verified for example by the aluminium depth profile, varies strongly from sample to sample. The measurements and the overall high oxygen content indicate that the sample surfaces may be very complex or even porous, This was partially confirmed by microphotography, see Fig. 4. The surface porosity of annealed Ai-Li alloys has also been reported by Papazian et al, [29]. This is one of the reasons why wall defined oxidation kinetics could not be measured, in contrast with successful results for titanium foils [28, 30].

TABLE I, Results of RBS and N D P analysis of AI-- Li alloy foils, annealed in prevacuum (typically 15 71 × I0 : "li~rrl Annealing.

Total ,x~ncentration

Temperature ( CI

Time (h : rain)

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D, Fink et al, / Oxidation of AI-Li alloys

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3.4, Annealing at ambient air pressure For this study, only the 100 pm thick AI-Li foils were used on account of the high oxidation speed, Owing to the high oxygen abundance at normal air pressure, all the samples looked thoroughly oxidized and the surface was roughened after annealing. Therefore, RBS measurements could not be performed with these samples, and our studies are restricted to NDP examinations. Thus, our NDP examinations do not yield real onedimensional lithium depth distributions, but rather the average over a variety of different surface structures. In contrast to the above findings at high vacuum, the annealing of the AI-Li alloy at ambient air pressure proceeded relatively fast, Figure 5 gives an overview of the lithium profile results of isochronal annealing (in 50 °C steps for I day each), and Fig. 6 gives an overview of the results of isothermal annealing for quadratically inca~asing times of approximately I, 2, 4.... h at 400 °C, Similar results are obtained at higher temperatures (450550 °C). Though the total annealing times were extended up to 3 months, it turned out that, within the depth of interest in our measurements (101am), no significant changes were observed after roughly 100 h annealing time. The general impression from the spectra is that they change very smoothly with annealing time and temperature, and that the depth profiles are very broad, extending much deeper than the accessible depth interval of about 10 I~m. (Our evaluation procedure used the triton spectra for the bulk regime (down to more than 10 pm depth), and the = spectra for details of the surface (down to a few I~m depth).) These results show the following tendencies. lsochronal annealing. Lithium segregates very rapidly at the surface, where its concentration remains roughly constant at successively higher annealing temperatures. The subsequent layer becomes depleted of lithium, the width of the depletion zone increasing with increasing temperature. Finally, the depletion width exceeds the measurable depth interval, so that no further profile changes are observed. When the temperature exceeds 500 °C, lithium becomes slightly depicted at the surface, This may be due to the onset of sublimation.

350

Fig. 3. (a) Two examples of oxygen depth distributions in annealed AI-Li alloys, as measured by RBS: O, annealing at 10 -2 Tort for 16 rain at 500 ~C (5th line of Table i); O, annealing for 40 rain at 500 "C, same pressure (line 6 in Table I). For comparison, the distribution of oxygen in a calibration sample (120nm SiOz on silicon) is shown (smooth line). (b) Comparison of lithium and oxygen depth profiles in the surface area of an AI-Li alloy foil, annealed at 10 -a Tort, 500 ~'C for 1632 h. Stoiehiometry is not reached at the surface, but somewhat below.

Isothermal annealing, For all three temperatures, lithium first starts to precipitate rapidly at the surface (I h), The precipitation is strong at lower temperatures (400, 500 °C) and weak for higher temperatures (550 °C). A subsequent deeper zone (about 6 pm thick after l h annealing at 400 °C) is always depleted of lithium, before the lithium concentration increases again at greater depths. After prolonged annealing, the broadening of the depleted zone exceeds the measurable depth interval of about l0 I~m. The final lithium distribution is in all

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D, Fink et al, / Oxidation oJ"AI-Li alloys

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cases rather homogeneous, with a slight increase in concentration towards the surface. AES examinations of commercial AI-Li-Mg alloys indicate that, after oxidation at ambient air, magnesium also behaves similarly to IRhium, by co-precipitating at the surface [20]. 3.5. Contaminations The RBS analysis did not only show the expected system components aluminium, lithium and oxygen, but (for the saml~les annealed in prevacuum) an additional contamination of superficial carbon, also listed in Table l. The carbon contamination is astonishingly high. It probably arises from thermal cracking of the oil vapour (from the prevacuum pump) and eventually of residual traces of acetone at the sample surface. After annealing for about 8 min, the concentration of carbon is fairly constant, independent of further annealing. At present, it is unknown whether this carbon contamination affects the oxidation kinetics to some extent. Strong carbon signals were also found by AES for untreated commercial AI-Li-Mg alloys and pure aluminium foils [20]. The shape of the Auger emission lines suggested a graphite form of carbon impurity. Our NDP technique is capable of observing l°B simultaneously with 6Li, via the l°~n,~t)~Li reactions, see refs. 24-26. Our measured NDP spectra show, in addition to the lithium depth profiles, a transient precipitation of the intrinsic boron impurity (typically of the order of some ppm atomic), which starts around ~ °C, and reaches a maximum after 32 h at 400 °C (or after 8 h at 500 +C and after 4 h at 550 °C). At longer annealing times or at higher temperatures, this boron precipitation dissolves again.

4. Ccmdmioas The oxidation process of AI-Li alloys exhibits a very complex behaviour. Below 350 °C. the foil composition does not change, which indicates very low lithium mobility, in contrast to earlier diffusion measureanents with lithium-implanted aluminium foils. Above this temperature, lithium moves out of the bulk towards the surface, where it is trapped as oxide. It appears that, during annealing in vacuum, the oxidation initially proceeds very fast but subsequently hardly changes. After short annealing times, the oxide is not yet stoichiometric> but Li20 stoichiome.try is reached after prolonged annealing. The oxygen and lithium distributions of the surface oxide consist of a nearsurface peak and a tail with decreasing concentration. extending into the bulk. SEM images show the evolution of an inhomogeagous surface with local oxide clustering. Also, previous reports on superficial pore formation are reconfirmed in this study. Below this surface oxide hyer, there exists a lithiumdepleted zone, which may extend far into the bulk. The thickness of this zone may he considerable, literature values range up to 0.3 mm. The present measurements with NDP show a gradual increase in the depletion zone, until the width exceeds the measurable depth interval, which is of the order of some 10 gm.

Ackw~gam~ This cooperation became possible by the recent political changes in Eastern Central Europe. One of the authors (D. F.) acknowledges the kind invitation by the

64

D. Fink el al. : O.
Institute of Nuclear Physics, I~eL and V, H, is grateful to the DFG for sponsoring his stay in Berlin, Further, we express our thanks for the support from the ILL Grenoble, where the N D P formed,

14

measurements were per15

Referexces I H. R6mpp. ('heroic der Metulle. Franckh'sche Verlagsbuchhandlung. W. Keller. Stutgart. 1949. p. 77. 2 E. S. Balmuth and R. Schmidt. in T. H. Sanders and E. A. Starke Icds.). Aluminium-Lithium AII,.L~. TMS;AIME. Warrendale. PA. 1981. p. 69. 3 A. M. Scnia. iron Age. 226 119831 47. 4 K. K. Sankaran. PhD "l'he.~is. Ma~sachusets Institute of Technology. Cambridge. MA. 1978. 5 C. Baker. P. J. Grcgson. S. J. Harris and C. J. Peel (~s.k Aluminium Lithium AIh~ys !11. Proc. 3rd Int. Aluminium I.ithimn Co~!l:. O.x:lord. July 8 II. 198.';. by Institute of Metals. London. 1986. G. Champier. B. du l~sI. D. Miannay and L. Sabctay (eds.). Alundnium Lithium AIIoy.~ IV. Proc. 4th h~t. .41,minium Lithium ('o~!/~. Puri.~. June IO 12. 1987. in 3. Ph)'s. ~'Puri.~. 43. ('~dl. ('3. Isuppl. 91. Les editions de physique. Les Ulis. 1987. 6 A. K. Fisher and D. R. Vis.~crs. 3. Electruchem. Sot.. l.ffl (19831 5. 7 D. M. Chen and H. F. Giddart. J. Electr~when,. Sot.. 135 (1983) 1975. 8 Z. Tomczuk. L. Redey and D. R. Visors. 3. Electr~a'hem. Sot'.. 13¢t 119831 1074. 9 N. P. Yao. L. A. Hercdy and R. C. Saunders. 3. Electrochem. S,c.. Ih~¢ 119711 1039. I(I E. J. Cairns and R. K. Steunenberg. in C. A. Rouse ted.). Pro.~rex~ in ll(t~h Temperature Phyxic.~ and ChemL~try. Vol. 5. Pergamon. New York. 1973. p. 63. II E. C. Gay. D. R. Vissers. F. J. Martino and K. E. Anderson. 3. Electrochem. S~'.. 123 (1976) 1591. 12 J. E. Battles. in C. Stein led,), Criti('ul Muteriul.~ Pr,hlem:~ i. En.~ineerbk~ Pr~alucti,n. Academic Pre~. New York. 1976. p. 769. 13 R. K. Steunenberg and M. F. Roche. in J. D. E. Mclntyre. S. Srinivasan and F. G. Will (eds.). Proc. Syrup. on Electrode Muteriuls

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