Joseph Anthony Smerdon

Julian Ledieu

Jon T Hoeft

Dennis E Reid

Lisa H Wearing

Renee D Diehl

Ronan Mcgrath

Thomas A Lograsso

Amy R Ross

Adsorption of cobalt on the ten-fold surface of d-Al72Ni11Co17 and on the five-fold surface of i-Al70Pd21Mn9

Keywords: scanning probe microscopy, quasicrystals, diffraction pseudomorphic PACS numbers: 61.44.Br, 68.35.Bs, 68.37.Ef

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Quasicrystals are complex alloys, usually consisting of three or more elements arranged in a highly ordered but aperiodic fashion [1]. The effect of aperiodic ordering on the physical properties of these exotic materials is inextricably linked to those properties that arise from the comparatively large number of species present. There is therefore much interest in single-element aperiodic structures, the formation of which may be facilitated by using the aperiodic surfaces of quasicrystals as nanotemplates for the adsorption of single-element films [2][3][4][5][6][7].

In this paper we report the formation of a pseudomorphic Co layer on the ten-fold surface of d -Al 72 Ni 11 Co 17 at 300 K; the formation of a less well-ordered Co thin film on the five-fold surface of i -Al 70 Pd 21 Mn 9 is also reported. The quasiperiodic structure of the films is evidenced in low-energy electron diffraction (LEED) patterns, and in part from scanning tunneling microscopy (STM) measurements taken from the films.

II. EXPERIMENTAL DETAILS

The d -Al 72 Ni 11 Co 17 sample was produced in the Ames laboratory using the melt decantation method, and the i -Al 70 Pd 21 Mn 9 sample by the same group using the Bridgman method. The ex-situ and subsequent in-situ preparation procedures are described elsewhere [8].

Using an Omicron EFM-3 e-beam evaporator, Co was deposited onto the ten-fold d -Al 72 Ni 11 Co 17 surface at a rate of 6×10 -3 ML sec -1 and onto the five-fold surface of i -Al 70 Pd 21 Mn 9 at the same rate (as determined by STM coverage measurements and by a comparison of the ion flux as measured by the evaporator). All deposition was carried out with the substrates at a temperature of 300 K. Scanning tunneling microscopy, LEED and AES measurements were taken from the resulting films. The error on the coverages quoted for Co/d -Al 72 Ni 11 Co 17 is estimated at 5% and for Co/i -Al 70 Pd 21 Mn 9 it is estimated at 20%. At higher coverages (5.6 ML and above) evidence for a row structure of the film is observed. It can be seen from the emergence of straight line domain boundaries in the STM images shown in figure 1 that certain orientations are preferred for the formation of islands of Co. At higher coverages, some one-dimensional row structuring within these islands becomes apparent. Due to limitations in resolution, it cannot be shown conclusively from STM that the inter-row spacings are arranged in a Fibonacci sequence, as we have previously observed for Cu adsorption on the five-fold i -Al 70 Pd 21 Mn 9 surface. The step height of the film as determined from STM is 2.0 ± 0.1 Å, which is consistent with the interlayer separation of hcp Co(0001) previously observed with STM to be 1.95 ± 0.10 Å [9]. The fact that the Al/Co Auger peak ratio drops to zero with increasing coverage implies that there is no alloying within the film.

III. RESULTS

A. Scanning tunneling microscopy and Auger electron spectroscopy

Scanning tunneling microscopy images taken from the Co/i -Al 70 Pd 21 Mn 9 film revealed no definite order, at any coverage (determined by STM and checked with AES), though at higher coverages there is a suggestion of linear structures within the incomplete layers formed. Auger electron spectroscopy again implies that there is no alloying within the film. Figure 2 (a,b) indicate that the dark five-fold hollows are not a preferred nucleation site for this adsorbate (as they have been found to be in previous studies [10]) as they may still be clearly observed when there is a coverage of 0.4 ML Co on the surface, with a density of approximately half that found on the clean surface. The size of the five-fold hollows (around 8 Å) indicates that they belong to the substrate.

B. Low-energy electron diffraction

The LEED patterns obtained from the multilayer films show more evidence for ordering. Figure 3 a periodic ordering along the Fibonacci rows, and a 2D LEED analysis was carried out to determine this periodicity.

Although we cannot directly observe Fibonacci ordering from STM of the Co/d -Al 72 Ni 11 Co 17 film, the similarity of the diffraction patterns leads us to conduct a similar investigation into this system. Using the PhD thesis of Sharma [START_REF] Sharma | Structure, Morphology, and Dynamics of Quasicrystal Surfaces[END_REF], a diffraction peak of known momentum transfer (k = 2.67 Å-1 ) indexed as (1221), was identified in the d -Al 72 Ni 11 Co 17 diffraction pattern shown in figure 4 (a). The momentum transfer k of the streaks in the Co/d -Al 72 Ni 11 Co 17 film was hence calculated from the ratio of the line length y/x using k = k.y/x and was found to be 2.56 Å-1 . Hence the periodic spacing along the rows is found to be 2π/k = 2.5 ± 0.1 Å, a figure which corresponds well to the nearest neighbour distance in the stable hcp phase of cobalt of 2.51 Å.

IV. DISCUSSION

On the five-fold surface of i -Al 70 Pd 21 Mn 9 , there is no suggestion of order in the STM images until a relatively high coverage of Co is reached and even then it is not well resolved. This order is also not apparent in diffraction patterns taken from the film until similar high coverages are reached. The growth as determined by STM is most closely approximated by the layer-by-layer model, though the layers are rough, and the island size small (tens of angstroms).

The Co/i -Al 70 Pd 21 Mn 9 film produces a diffraction pattern that, within the error introduced by the correction for optics, and indeed the quality of the patterns themselves, is extremely similar to that from the Cu/i -Al 70 Pd 21 Mn 9 film. Assuming the streaks observed in the pattern from the Co film indicate periodicity, as they do for the Cu film, then we may conclude that they correspond to the same periodicity, as far as is permitted within the accuracy of such an analysis. The NN distance in bulk fcc Cu is 2.56 Å. Our LEED analyses of the Cu film and of the Co films yield a periodicity along the rows of 2.5 ± 0.1 Å. We have used this analysis to assist our assertion elsewhere [START_REF] Ledieu | [END_REF] that the structure of the Cu phase present on i -Al 70 Pd 21 Mn 9 consists of five orientational domains composed of rows of atoms arranged according to a one-dimensional Fibonacci sequence.

On the ten-fold surface of d -Al 72 Ni 11 Co 17 , deposited Co appears to form an initial complete layer upon which smaller domains of row-structured material take shape. With each subsequent layer, the lateral island size decreases. The growth is therefore neither strictly layer-by-layer nor 3D, though layer-by-layer growth is the closest approximation. At the highest coverages studied (15ML), the size of the Co islands was on the order of hundreds of angstroms. The streaks present in the diffraction pattern from the Co/d -Al 72 Ni 11 Co 17 film can immediately be seen to correspond to a smaller period (due to the significantly larger separation of the streaks on the pattern). The NN distance in bulk hcp Co is 2.51 Å. Our LEED analysis of the Co/d -Al 72 Ni 11 Co 17 film yields a periodicity along the rows of 2.5 ± 0.1 Å. The momentum transfer extracted for the separation of the streaks on the Cu/i -Al 70 Pd 21 Mn 9 was 2.48 Å-1 [START_REF] Ledieu | [END_REF]. To the extent that the values of the momentum transfers derived from the bulk structures are correct for these surfaces, we can see that the difference in the momentum transfers for the two films is (2.56/2.48 = 1.03) and this translates to a difference in lattice parameter of 3%. We can thus determine that the periodicity for Co on i -Al 70 Pd 21 Mn 9 is 3% larger, in one dimension, than that on d -Al 72 Ni 11 Co 17 . Bulk fcc Cu has a NN distance 2% larger than that of bulk hcp Co. The table in figure 5 lists these properties.

V. CONCLUSIONS

We find that, when deposited onto the ten-fold surface of d -Al 72 Ni 11 Co 17 , Co grows in a predominantly layer-by layer fashion, forming a film with a number of orientational domains. These domains have a row structure visible in STM images, and are shown by a LEED analysis to be consistent with a one-dimensional Fibonacci arrangement of rows, with a periodic arrangement of atoms perpendicular to the Fibonacci direction. The NN distance along these rows has been determined to be 2.5 ± 0.1 Å, and has also been shown to be smaller than the periodicity of the Cu/i -Al 70 Pd 21 Mn 9 structure by around 3%. This is consistent with the NN distance in bulk hcp and fcc Co; however, due to the interlayer spacing observed by STM, it is suggested that the structure adopted by the Co on this surface is composed of rows of hcp structured material with the (1000) hcp orientation along the rows. It has also been observed that a film of Co deposited on i -Al 70 Pd 21 Mn 9 appears to have the same broad structure and periodicity as a film of Cu deposited onto the same surface. However, the Co film is far less ordered than the Cu film on this surface, and also less ordered than the Co film deposited on d -Al 72 Ni 11 Co 17 .

In the regime of pseudomorphic growth, there have been to date three multilayer films formed on quasicrystal substrates: the copper film on i -Al 70 Pd 21 Mn 9 [7] and the studies under discussion. Although the structural properties of the two substrates are very different (i -Al 70 Pd 21 Mn 9 is an icosahedral quasicrystal with three-dimensional aperiodicity, whereas d -Al 72 Ni 11 Co 17 is periodic along one direction and decagonal and aperiodically ordered perpendicular to that direction), common features emerge in the films deposited: the atoms order themselves in row structures according to orientations present within the substrate, and these row structures are themselves ordered according to a one-dimensional Fibonacci sequence.

  Adsorption of cobalt on the ten-fold surface of d -Al 72 Ni 11 Co 17 and on the five-fold surface of i -Al 70 Pd 21 Mn 9 I. INTRODUCTION

2 FIG. 1 :

 21 FIG.1: (a-g); Comparison of the ten-fold d -Al72Ni11Co17 surface as studied with increasing Co coverage. The coverage is respectively 0.4 ML, 1.9 ML, 3.7 ML, 5.6 ML and 7.5 ML (for (e-g). (f ) is a detail of (e), with the row structure showing more clearly. Indications of orientational order are superimposed on (e-g). A graph showing the trend of Auger peak area ratios vs coverage is also shown.

Figure 1

 1 Figure1shows images taken from the Co/d -Al 72 Ni 11 Co 17 film at room temperature with an Omicron variabletemperature STM. No order is discernible at 0.4 ML coverage (determined by STM and checked with AES).At higher coverages (5.6 ML and above) evidence for a row structure of the film is observed. It can be seen from the emergence of straight line domain boundaries in the STM images shown in figure1that certain orientations are preferred for the formation of islands of Co. At higher coverages, some one-dimensional row structuring within these islands becomes apparent. Due to limitations in resolution, it cannot be shown conclusively from STM that the inter-row spacings are arranged in a Fibonacci sequence, as we have previously observed for Cu adsorption on the five-fold i -Al 70 Pd 21 Mn 9 surface. The step height of the film as determined from STM is 2.0 ± 0.1 Å, which is consistent with the interlayer separation of hcp Co(0001) previously observed with STM to be 1.95 ± 0.10 Å[9]. The fact that the Al/Co Auger peak ratio drops to zero with increasing coverage implies that there is no alloying within the film.Scanning tunneling microscopy images taken from the Co/i -Al 70 Pd 21 Mn 9 film revealed no definite order, at any coverage (determined by STM and checked with AES), though at higher coverages there is a suggestion of linear structures within the incomplete layers formed. Auger electron spectroscopy again implies that there is no alloying within the film. Figure2 (a,b) indicate that the dark five-fold hollows are not a preferred nucleation site for this adsorbate (as they have been found to be in previous studies[10]) as they may still be clearly observed when there is a coverage of 0.4 ML Co on the surface, with a density of approximately half that found on the clean surface. The size of the five-fold hollows (around 8 Å) indicates that they belong to the substrate.

3 FIG. 2 :

 32 FIG. 2: (a-d); Comparison of the five-fold surface of i-Al70Pd21Mn9 as studied with increasing Co coverage. is (a) 0.4 ML; (b) as (a) with five-fold depressions highlighted; (c) 2.5 ML coverage, and (d) 13 ML coverage.

  shows the patterns obtained at 2 different electron energies from the Co/i -Al 70 Pd 21 Mn 9 film, the Co/d -Al 72 Ni 11 Co 17 film (both at room temperature) and previously obtained patterns from the pseudomorphic Cu/i -Al 70 Pd 21 Mn 9 film[7] (at 85 K). The patterns on the right have been corrected for the different acceptance angle and size of the optics so they may be compared with the others. Each coverage quoted provided the optimum LEED pattern in terms of lowest background and clearest diffraction peaks. It can be observed, despite the lower quality of the patterns, that the Co/i -Al 70 Pd 21 Mn 9 diffraction is qualitatively similar to that for Cu/i -Al 70 Pd 21 Mn 9 , and the Co/d -Al 72 Ni 11 Co 17 patterns, apart from having a different set of momentum transfers, show the same broad characteristics as both.In diffraction obtained from the Cu/i -Al 70 Pd 21 Mn 9 film, the periodic arrangement of the streaks was attributed to http://mc.

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manuscriptcentral.com/pm-pml Philosophical Magazine & Philosophical Magazine Letters

  LEED pattern from a 7.5 ML Co overlayer at 120 eV. Lines are superimposed to aid identification of the periodic separation of the streaks.

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	16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 F o r Lattice NN distance parameter ( Å) ( Å) hcp Co --2.51 fcc Co 3.56 2.52 P e r R e v i e w O n l y Interlayer Lattice parameter Lattice parameter 1.95 2.5(3) 2.4(6) 1.78 (001) ----bcc Co 2.82 2.44 1.41 (110) ----e fcc Cu 3.62 2.56 1.81 (001) 2.5(3) --

4 FIG. 4: (a); LEED pattern obtained from the clean surface at 120 eV. (b); spacing ( Å) (STM) on i-AlPdMn ( Å) on d -AlNiCo ( Å) FIG. 5: Table

showing

certain structural properties of copper and cobalt

[START_REF] Chambers | [END_REF]14]

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VI. ACKNOWLEDGEMENTS

This material is based upon work supported by the EPRSC under Grant No. GR/S19080/01, the National Science Foundation under Grant No. 0208520, the U. S. Department of Energy and the EU Marie Curie Host Development fund.


S Francoual
email: sonia.francoual@ltpcm.inpg.fr

F Livet

M De Boissieu

F Yakhou

F Bley

A Létoublon ¶

R Caudron||&

J Gastaldi#

ENSEEG R Currat ‡ †ltpcm

Dynamics of long-wavelength phason fluctuations in the i-Al-Pd-Mn quasicrystal

Keywords: X-ray, quasicrystals user supplied): phasons, coherent X-ray Quasicrystals, Phasons, Coherent X-ray

We report on the dynamics of phason modes in the i-Al-Pd-Mn icosahedral quasicrystal, measured between room temperature and 650°C, using the Xray Intensity Fluctuation Spectroscopy technique (XIFS). Up to 500°C, the auto-correlation function ℑ(q, t) displays almost no time evolution as expected for frozen-in phason fluctuations at low temperature. At higher temperatures, ℑ(q, t) follows a single exponential time decay from which the characteristic time τ c (q) is extracted. These results are compared with the expected shape of ℑ(q, t) as derived from the expressions of the eigenvalues and eigenvectors of the C ⊥⊥ (q) phason dynamical matrix. In agreement with the hydrodynamic theory of quasicrystals which predicts phasons with diffusive character, we find that τ c (q) varies linearly with q -2 at 650°C. The corresponding diffusion coefficient is 2.2(± 0.5)×10 -18 m².s -1 and the activation energy is estimated around 2.3(± 1) eV.

§1. Introduction

Quasicrystals are long-range ordered structures having rotational symmetries incompatible with translational invariance (see [1] and refs. therein). In the superspace description of icosahedral phases, the periodicity is restored in a 6-dimensional (6D) space, E. There, the lattice is a 6D cube decorated with 3D atomic domains lying in E ⊥, the perpendicular 3D space. E ⊥ is the phason space, oriented in such a way as to be orthogonal to the physical space E // and incommensurate with E . A section of E by E // generates the 3D atomic structure. Hydrodynamics is concerned with those excitations whose frequency ω → 0 as the wavevector q → 0. The symmetries leaving the system's free energy invariant are associated with u and w, the two uniform displacement fields along E // and E ⊥ , respectively [START_REF] Lubensky | Introduction to Quasicrystals[END_REF]. The degeneracy with respect to the variable w introduces 3 Goldstone phason modes in addition to the 3 acoustic phonons associated with u. Long-wavelength phasons are diffusive modes associated with discrete atomic rearrangements in physical space. In a simplified picture, a phason can be seen as a damped sine-wave fluctuation with wavevector q along E // and polarisation vector e ⊥ (q) along E ⊥ . Neglecting phononphason interactions, 3 phason eigenmodes (k = 1, 2, 3) can be defined for a given vector q, each having its own relaxation time τ ⊥, k (q) and coefficient of diffusion D ⊥, k ) ( q , given by :

2 , k , k 1 , k q ) ~ ( D ) ( K Γ ) ) ( τ ( × = × = ⊥ ⊥ - ⊥ q q q w (1)
where the subscript ⊥ stands for the phason related variables. Γ w is a kinetic coefficient and ~ q = q/|q|. K ⊥, k (q)and e ⊥, k (q) are, respectively, the eigenvalues and eigenvectors of the 3×3 C ⊥⊥ phason dynamical matrix. The C ⊥⊥ matrix elements depend on K 1 and K 2 , the phason generalised elastic constants, and on the components of q.

Long-wavelength phasons lead to additional diffuse scattering near Bragg peaks in reciprocal space [START_REF] Jaric | [END_REF][4][5]. In the framework of elasticity theory, generalised to icosahedral phases, the phason diffuse intensity, noted I ⊥ , at a distance q from a Bragg peak with component Q // in reciprocal space writes as :

∑ ∑ ⊥ ⊥ ⊥ ⊥ ⊥ ⋅ × × × = + = + k k , 2 k , B // Bragg k // k , // ) ( /K ) ) ( ( T k ) ( I ) ( I ) ( I q Q q e Q q Q q Q (2)
where Q ⊥ is the conjugate component of Q // in perpendicular space, k B the Boltzmann constant and T the temperature. Neutron and X-ray scattering experiments performed on i-Al-Pd-Mn have evidenced such a phason diffuse intensity response [6][7]. In particular, insitu studies show a decrease of the diffuse intensity, I ⊥ , simultaneously with an increase of the Bragg peak intensity, I Bragg , as the temperature is increased from 600°C to 770°C. This reversible temperature evolution is found related to a softening of the phason elastic constant K 1 , the ratio K 2 /K 1 changing from a value -0.53 below 600°C to a value -0.40 at 770°C.

In the present paper, we report on the dynamics of phasons having a wavevector q along the (-τ, 1, 0) direction, where τ is the golden mean defined as τ = (1+√5)/2. The diffuse intensity is measured using a coherent X-ray beam so that it exhibits sharp intensity fluctuations called speckles. Speckles are sensitive to the time evolution of the atomic density so that any time dependence induced by phason fluctuations will result in a correlated time dependence for the speckle intensity distribution. The intensity-intensity correlation function ℑ(q, t) is defined as:

, t) g( β 1 ) , t') I( /( t) ' , t I( ) ' t , I( , t) ( 2 t' ' t q q q q q × + = > < > + × < = ℑ (3)
with β the degree of coherence of the beam and g(q, t) the function which characterises the time dependence of the fluctuation. For a diffusive process, g(q, t) = exp(-2t/τ c (q))

where τ c (q) is a characteristic time † . §2. Experimental

The coherent X-ray scattering experiment was carried out on the ID20 beam line at the ESRF in Grenoble following the set-up described in [8,9]. The incoming energy is set to 8 keV, with an energy resolution ∆λ/λ = 1.3×10 -4 , yielding a longitudinal coherence length (λ²/∆λ) ~ 1.2 µm. A 10 µm collimating pinhole before the sample allows for a ~ 10 µm transverse coherence length together with a 5×10 8 photons/s flux at the sample position. Diffraction patterns were recorded on a Princeton Instrument Direct Illumination 2D CCD acting as a 2D X-ray photon counter [10]. About 500 frames were recorded at each temperature with a time interval of a few seconds (<10 s) between successive frames. The sample is a 1mm thick Al 68.2 Pd 22.8 Mn 9 plate-shaped single-grain with a 5fold axis perpendicular to its surface, extracted from a Czochralski grown mother ingot [11]. It was annealed for 8 days at 780°C under ultrahigh vacuum and then slowly cooled down in order to guaranty the surface stability during the X-ray measurements [12]. After annealing, the sample surface displays flat terraces whose size, between 50 and 100 µm, is sufficient to focus the 10 µm beam on it. The sample is placed in a furnace and heated up under secondary vacuum.

Measurements were performed near the 7/11 reflection having a Bragg angle 2θ

~ 24° and a component Q // along the (1, τ, 0) 5-fold axis. The partial coherence of the beam β is 5%. From room temperature up to 650°C, several q values were probed by rotating the θ incident angle with a 0.1° step which corresponds to a step q ~ 0.003 Å -1 along the (-τ, 1, 0) direction in reciprocal space. §3. Results

Figure 1 shows a typical diffraction pattern obtained at a position q = 0.0125 Å -1 away from the 7/11 Bragg peak. The simulation of the diffuse intensity distribution, shown on the right panel, is in good agreement with the measurement, insuring that the observed diffuse signal is due to phason modes. In order to improve statistics, ℑ(q, t) must be averaged over a set of equivalent q vectors. Given the anisotropy of the phason diffuse intensity distribution, we considered only an average performed over a circle close to the q //(-τ, 1, 0) wavevector at the center of the CCD. '[insert figure 1 about here]'

From room temperature up to 500°C, ℑ(q, t) is found almost constant over 1000s, as shown in figures 2a and 2b. The temperature range between 500°C and 600°C is a transient regime in which some dynamics sets up. The time evolution is however very slow as can be seen in Fig. 2c and only few data were thus recorded. At 650°C, ℑ(q, t) displays a clear exponential time decay from which the characteristic time τ c is extracted

(see Figs. 2(d-f)).

'[insert figure 2 about here]' '[insert figure 3 about here]'

The observed evolutions can be related to the dynamics of phason modes. Below 500°C, the absence of dynamics would be the signature of frozen-in phason fluctuations as reported in [7][8]. The temperature activation occurs between 500°C and 600°C. Above 600°C, the time evolution of ℑ(q, t) is exponential as expected for a relaxation process. Fig. 3 is a plot of the decay rate, (1/τ c ), as a function of q² at 650°C In agreement with equation (1), the dependence is linear which demonstrates that phason modes have diffusive character. The corresponding diffusion coefficient at 650°C is D ⊥ ) ~ ( q = 2.2(± 0.5)×10 -18 m².s -1 ‡ . Between 600°C and 650°C, τ c is divided by a factor ~ 5 at same q value (see Figs. 3c and3e). This leads to an activation energy E a = 2.3(±1) eV, assuming an Arrhenius-type behaviour. §4. Discussion and conclusions From hydrodynamic theory, for each q, one expects to observe three phason modes, each with its own relaxation time. The behaviour of the function ℑ(q, t) should thus reflect the existence of these three relaxation times through a more complex time evolution than the single exponential time decay that is experimentally observed. By considering the intensity and relaxation time of each mode, we demonstrate, in the following, that in fact, for this particular phason wavevector direction, two modes are selected and that the resulting time evolution of ℑ(q, t) is close to a single-exponential decay. ‡ The value of D ⊥ ) ~ ( q differs numerically by a factor 8π² from the value quoted in ref [8]. The diffuse intensity is measured at the position (Q // +q) where q is along (-τ, 1, 0) and Q // along (1, τ, 0). In that direction of q, there exists a basis e ⊥, k (q) where C ⊥⊥ (q) is diagonal with eigenvalues K ⊥, k (q):

q //(-τ, 1, 0) →               × × - = × × + - = - × × - + = ⊥ ⊥ ⊥ ⊥ ⊥ ⊥ ) 1 , 0 , 0 ( // ) ( , q ] K ) 3 / 1 ( K [ ) ( K ) 0 , 2 / 1 , 1 ( // ) ( , q ] K ) 3 / 1 /τ 1 ( K [ ) ( K ) 0 , 1 , 2 / 1 ( // ) ( , q ] K ) 3 / 1 (τ K [ ) ( K 3 , 2 2 1 3 , 2 , 2 2 1 2 , 1 , 2 2 1 1
, q e q q e q q e q (4)

The phason diffuse scattering intensity is the sum of the three contributions I ⊥, k as expressed in equation ( 2). The scalar products

(Q ⊥ ⊥ ⊥ ⊥ . e ⊥, k ) act as selection rules. Since Q ⊥ ⊥ ⊥ ⊥ lies along (-τ, 1, 0) in E ⊥ space, it follows that (Q ⊥ ⊥ ⊥
⊥ . e ⊥, 3 ) is zero and the mode 3 does not contribute to the intensity. We can estimate the weight of the 1 st and 2 nd modes from the value of K 2 /K 1 . We find I ⊥, 1 ~ I ⊥, 2 at 770°C and I ⊥, 1 ~ 2I ⊥, 2 below 600°C. We thus select two modes which contribute with comparable weights to the overall diffuse intensity at all temperatures.

The time evolution of ℑ(q, t) depends on the time evolution of the two selected phason modes. More precisely, in definition (3), g(q, t) writes as [START_REF] Berne | Dynamic Light Scattering[END_REF] :

2 , 2 , 2 , 1 , 1 ] )) ( t/τ exp( ) /I (I )) ( t/τ exp( ) /I (I [ , t) g( q q q ⊥ ⊥ ⊥ ⊥ ⊥ ⊥ - × + - × = (5)
The phason relaxation times τ ⊥, 1 and τ ⊥, 2 depend on Γ w and K ⊥, k (see equation ( 1)). Both parameters have their own temperature behaviour. In the diffuse scattering measurements, we probe the temperature dependence of the K 1 and K 2 elastic constants only, whereas in the present experiment, we measure directly τ ⊥ and thus get the complete temperature information. Assuming that the kinetic coefficient Γ w is identical for the three phasons, we can estimate that τ ⊥, 1 ~ 5τ ⊥, 2 at low temperature while τ ⊥, 1 ~ 3τ ⊥, 2 at 770°C. Given the intensity and decay time ratios obtained between the 1 st and 2 nd phason modes, numerical simulations show that, within the statistical accuracy of our data, it is not possible to distinguish the time evolution given by equation ( 5) from a pure single exponential decay.

The characteristic decay time τ c (q) that is then extracted is an effective relaxation time averaged over the two selected phason modes. At 650°C, the phason diffusion coefficient D ⊥ ) ~ ( q is found equal to 2.2×10 -18 m².s -1 . When compared with the atomic self-diffusion of the Pd and Mn elements obtained at same temperature in i-Al-Pd-Mn by the radiotracer techniques [START_REF] Mehrer | Diffusion in Quasicrystals, Quasicrystals -Structure and Physical Properties[END_REF], we find 4 orders of magnitude of difference between the two types of relaxation processes, as D Pd ~ 4×D Mn ∼ 6×10 -14 m².s -1 . On the other hand, our result compares well with mechanical spectroscopy measurements [START_REF] Damson | [END_REF] in which a slow relaxation process involving the collective motion of a large number of atoms was identified. Phasons are believed to play a key role in the high temperature mechanical properties of quasicrystals. In particular, it has been observed that the motion of dislocations in i-Al-Pd-Mn [16] is followed by the formation of planar faults associated with the phason-displacement field w around the dislocation line. A possible scenario is that these phason walls disappear as phason modes are activated at high temperature. Such annealing has been recently reported in [17]. The complete dispersion of a ~0.1 µm thick phason wall is achieved after 12s at 700°C. Assuming that the above time scale is of the order of τ ⊥ , from the values of E a and D ⊥ , we obtain a characteristic phason length scale λ ⊥ (=2π/q) of 685 Å at 700°C. This is of the same order of magnitude as the phason wall thickness which indicates that both results are compatible.

To summarise, we measured the dynamics of long-wavelength phasons in the i-AlPdMn quasicrystal using coherent X-rays. The diffusive character of phason modes, predicted by hydrodynamic theory, is experimentally evidenced. 1: Left: Diffraction pattern obtained at 650°C at the distance q = 0.0125 Å -1 from the 7/11 Bragg peak. 400 frames were measured with a 7s time interval. The diffuse intensity distribution is anisotropic in agreement with an intensity response due to phasons. Right: Simulation of the phason diffuse scattering in the i-Al-Pd-Mn phase in the relevant part of reciprocal space following the model of Jaric [5] with K 2 /K 1 = -0.53. Figure 2: Time dependence of the auto-correlation function ℑ(q, t) at different temperatures and different values of the phason wavevector q. The solid line is an exponential fit to the data. ℑ(q, t) is constant up to 500°C (a-b). At 600°C and 650°C, ℑ(q, t) decays exponentially (c-f). At 650°C, the decay time τ c (q) decreases with q (d-f).

Figure 3: Evolution of the inverse of the relaxation time (1/τ c ) as a function of q² with q along (-τ, 1, 0), at 650°C. The linear trend is a signature of a diffusive process allowing the determination of the phason diffusion coefficient D ⊥ ) ~ ( q = 2.2(± 0.5)×10 -18 m².s -1 .

  Figure 1

The diffusion coefficient http://mc.manuscriptcentral.com/pm-pml Philosophical Magazine & Philosophical Magazine Letters

	is found to be 2.2(± 0.5)×10 -18 m².s -1 at 650°C with an activation energy equal to 2.3(± 1) Figure
	F o eV. F o
	r r
	P P
	e e r e e r
	R e v i e w R e v i e w
	O n O n
	l y l y

† In a previous paper[8], we overlooked the factor 2 which appears in the exponent.

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Javier Sebastián Blázquez

Victorino Franco

Alejandro Conde
email: conde@us.es

Stefan Roth

J S Blázquez

Prof A Conde

Ball milling nanocrystallization of arc-melted and melt-spun Fe 78 Co 5 Nb 3 Zr 3 B 5 Ge 5 Cu 1 alloy: Microstructure and magnetic properties

Keywords: nanocrystals, magnetic properties user supplied): mechanical alloying Nanocrystalline Materials, Mechanical Alloying, Magnetic Properties

published or not. The documents may come

Ball milling nanocrystallization of arc-melted and melt-spun Fe78Co5Nb3Zr3B5Ge5Cu1 alloy: Microstructure and magnetic properties

Introduction

Soft magnetic nanocrystalline alloys with a microstructure consisting of nano-sized α-Fe type crystallites embedded in a residual amorphous matrix are among the softest magnetic materials known up to date [1]. Generally, the microstructure of these alloys is achieved in a procedure consisting of two steps: first, an amorphous alloy is obtained with the desired composition and, afterwards, this precursor is conveniently annealed to develop a nanocrystalline microstructure. In order to obtain an amorphous alloy by quenching a melted system with a suitable composition to develop a nanocrystalline microstructure, it is necessary to use very high quenching rates (~10 5 -10 6 K/s). These very high cooling rates limit the geometry of the material to films or ribbons with a maximum thickness of few tens of microns. In order to overcome this limitation, consolidation of metallic powders has been proposed as a suitable technique to produce bulk shaped materials [2], widening their possible technological applications. In this sense, mechanical milling has been shown as an advantageous procedure to obtain nanocrystalline powder in many different composition systems [3,4]. However, the coercivity of the consolidated powder is 2-3 orders of magnitude larger than the values observed in melt-spun nanocrystalline ribbons, which has been associated to the absence of a residual amorphous matrix in the case of the mechanically milled nanocrystalline samples [2]. Therefore, the distribution of the atoms out of the α-Fe nanocrystals, forming intermetallics or segregated at the grain boundaries and its thermal stability are interesting subjects to study.

In this work, the preparation of a nanocrystalline Fe 78 Co 5 Nb 3 Zr 3 B 5 Ge 5 Cu 1 alloy by milling has been studied. The addition of Ge and Co to Fe-based nanocrystalline alloys has been proposed in order to enhance the magnetic properties of the system at [5][6][7]. There are a lot of parameters influencing the milling process for a given composition: type of mill, time, speed and media of milling, ball-to-powder ratio, starting material, etc. This study is devoted to the dependency of the milling process on the microstructure of the starting system for the same composition and, therefore, the other milling variables were kept constant. Two initial samples were chosen: an arc-melted ingot and a partially crystallized ribbon obtained by meltspinning. Both systems are compositionally homogeneous at the scale of the powder particles (~µm). This fact implies several advantages in comparison with the use of elemental powders as starting materials, especially in such complex compositions as the one studied in this work. The milling of elemental powders would need much more time to homogenize the composition and it would be quite hard to assure that the final composition will remain as the initial proportion of powder used due to preferential sticking of some elements on the milling media. On the other hand, a homogeneous composition is less reactive than pure elements, which will reduce the effect of oxygen contamination in the powder.

Experimental

A master alloy with the nominal composition Fe 78 Co 5 Nb 3 Zr 3 B 5 Ge 5 Cu 1 was obtained by arc-melting several times in an argon atmosphere and after purifying the atmosphere by melting titanium (AM sample). A portion of this alloy was melted again and rapidly quenched by using a single-roller Bühler melt-spinner (RQ sample). The ball milling was performed under argon atmosphere in a Retsch PM4000 planetary ball mill using hardened steel balls and bowls. The rotation speed was 150 rpm for all the experiments and the ball to powder ratio was 10:1. The RQ ribbon, ~100 µm thick, was very brittle and thus, it was easy to make flakes which were used as the starting material for the milling procedure. After 2 h of milling, these flakes are still present and thus it is not possible to consider this RQ sample as a real powder at this stage.

In order to obtain pieces of the AM sample small enough to start the milling in the planetary mill, big pieces (several mm) were pulverized using a SPEX shaker mill for short times (< 0.5 h). The XRD pattern of the obtained AM powder after milling in the SPEX did not show significant differences to the bulk AM alloy, except for a broadening of the diffraction maxima.

The thermal stability of the samples was studied by annealing the powder in an argon atmosphere, using a differential scanning calorimeter (DSC) Netzsch DSC 404.

The melting interval of the studied composition starts about T m ~ 1450 K, as shown in

Results and discussions

Microstructure

Figures 2 and3 show the XRD patterns of the RQ and AM samples after different times of milling. In the case of the AM sample labelled as-cast, the XRD pattern corresponds to the arc-melted ingot and, in the case of the RQ sample, it corresponds to the meltspun ribbon. As it can be observed for both as-cast samples XRD patterns present α-Fe crystalline peaks but also other diffraction maxima of some crystalline phases can be detected. The XRD patterns of the RQ samples show peaks corresponding to metastable phase(s), which become negligible after 5 h of milling. However, the diffraction maxima of the AM samples can be indexed using the peaks of NbFeB and Fe 2 Zr phases.

Unlike for RQ samples, these peaks do not disappear even after 70 h of milling. A tentative explanation of this fact must take into account the lower cooling rate of the melted composition in the case of the AM samples in comparison with that of the RQ samples. This provokes that the formation of stable crystalline phases different to α-Fe is overcome in the case of the initial RQ ribbon (quenching rate of ~10 5 -10 6 Ks -1 ), although some metastable crystalline phases appear. The metastability of these phases could be the responsible of the easiness for destroying them during milling, while the energy supplied by milling during our experiments was not enough to destroy the thermodynamically stable phases (α-Fe, NbFeB or Fe 2 Zr in the initial AM sample).

'[Insert figures 2 and 3 around here]'

The evolution of the α-Fe phase with the milling time and the thermal treatment was studied for both sets of samples after fitting the XRD profile in the 2θ range process. In the case of the RQ samples, the Lorentzian contribution is higher (decreasing from ~1 to 0.7 as the milling time increases, being ~1 after annealing). This indicates that the effect on the broadening of the (110) maximum for the RQ samples due to microstrain relative to small crystal size increases with the milling time, unlike for AM samples. From the broadening of the (110) deconvoluted maximum, it was possible to estimate the grain size, <D>, and the microstrain, ε. On the other hand, the lattice parameter was measured using the (110) and (200) diffraction maxima of the α-Fe phase. These results appear in table 1.

'[Insert table 1 around here]'

As it is expected, the microstrain of the samples increases with the milling time for both RQ and AM systems from 0.006 up to 0.020 (after 15 h of milling) and up to 0.008 (after 70 h of milling), respectively. This indicates a more significant and much faster increase of the microstrain for RQ samples than for AM samples. After annealing, the microstrain of the RQ samples could not be measured due to the negligible Gaussian The decrease in grain size of the α-Fe phase saturates in both types of samples at 8 nm, which is in good agreement with the value observed for pure Fe [8]. However, this value is obtained after 5 h of milling for the RQ alloy and only after 50 h of milling for the AM alloy. On the other hand, only the RQ samples milled during 10 and 15 h and heated up to 975 K remain with a value of D below 20 nm. For AM samples, the annealing process increases the grain size to 20 nm or higher, even for milling times of 70 h. This fact indicates that a certain microstructural stability can be achieved for the RQ sample after 10 h of milling but the microstructure of the AM samples is less thermally stable even after 70 h of milling. In fact, experiments performed on RQ sample heated up to 573 K show that the nanocrystal size increases from 8 to 15 nm for the alloy milled during 5 h, but the samples milled 10 and 15 h show the same crystal size after that thermal treatment. Figure 4 shows the XRD pattern of those RQ samples heated up to 973 K, for which α-Fe was the only crystalline phase detected after milling.

'[Insert figure 4 around here]'

The lattice parameter, a, does not differ for both as-cast samples (0.2874 and 0.2873 nm for the RQ and the AM alloys, respectively), but presents a different evolution during the milling and annealing processes. In the case of RQ samples, a is almost constant throughout the milling process and after annealing the sample.

However, for AM samples, the lattice parameter increases for long milling times up to 0.2884 nm and a clear reduction of a down to ~0.2860 nm is detected after annealing, even for samples milled during short times. Figure 5 shows the a values as a function of '[Insert figure 5 around here]'

The constant value of a for RQ samples is clearly larger than that of pure α-Fe (0.28663 nm) indicating the presence of some stable impurities in the crystal lattice.

Although long milling time of AM samples yields an increase of the lattice parameter of the α-Fe phase, it is worth noting that RQ samples were not submitted to such long milling times and, therefore, this increase could be ascribed to the higher energy accumulated for AM samples.

The decrease of a observed after annealing AM samples occurs even for short milling times, comparable to those of RQ samples, for which no change in a is detected.

This fact might be ascribed to the presence of stable intermetallic phases in the AM samples unlike for the RQ samples. The elements of the studied composition with better solubility in the α-Fe phase are Co and Ge. Although Co is highly soluble in α-Fe [START_REF] Massalski | Binary alloys phase diagrams[END_REF],

the solubility of Ge is much more reduced (up to 10 at % at room temperature) [START_REF] Massalski | Binary alloys phase diagrams[END_REF]. The mixing enthalpies of Ge with Nb and Zr (-24 and -25 kJ/mol, respectively) are much lower than those of Fe with the two other elements (-3 and -1 kJ/mol, for Nb and Zr respectively) [START_REF] De Boer | Cohesion in metals. Transition metal alloys[END_REF]. Therefore, it is possible that during annealing Ge atoms diffuse to the intermetallics, rich in Nb and Zr, and thus the α-Fe phase will reduce the concentration of this element, in agreement with the reduction of a observed for annealed AM samples. The value of a observed for RQ samples and not annealed AM samples after low milling times, ~0.2873 nm, would correspond to a Ge content of 4 at % [START_REF] Pearson | A handbook of lattice spacings and structures of metals and alloys[END_REF], as Co does not influence the lattice parameter of α-Fe below 10 at % and for higher content a decreases [START_REF] Bozorth | Ferromagnetism[END_REF]. The other elements Nb, Zr and B are strongly insoluble in the α-Fe phase. For AM samples, these elements might be found preferentially in the intermetallics; however, in the case of RQ samples, the boundaries between nanocrystals are suitable regions to be enriched in these elements.

Magnetic properties

In order to perform magnetic measurements, the powder samples were encapsulated in silver foil, which was folded in order to avoid the movement of the powder inside.

Saturation magnetization, M S , was measured in the VSM (the maximum applied field was 1500 kA/m) and the Koerzimat was used to obtain a value of coercivity, H C , after magnetizing the sample with an applied field of 200 kA/m. The resultant values appear in table 2.

'[Insert table 2 around here]'

It is worth noting that these values of H C were obtained in such a system consisting of a mixture of powder and air gaps. However, considering the repetitiveness of the measurement (low errors bar) and that all the samples were prepared in the same way, the reported values could be considered as self-comparable under this study. On the other hand, the RQ powder milled during 15 h and the AM powder milled during 70 h were pressed at 973 K during 5 min in order to form bulk disks (1 cm in diameter and several mm high) and the coercivity measured in these two bulk samples is in good agreement with those of the corresponding powder samples encapsulated in silver and heated up to 973 K (see table 2). Once this consideration is established, it is possible to observe lower values of coercivity for the AM samples than for the RQ samples, about 2500-3000 and 4500-5000 A/m, respectively. In both cases, annealing of the samples reduces the value of H C . Figure 6 shows the H C values for the different studied samples. It is worth noting that RQ sample milled during 2 h is not a real powder sample but consist of ribbon flakes. Therefore, this sample was discarded for the annealing.

'[Insert figure 6 around here]'

The M S values do not significantly change with the milling time or the annealing process. Although average values observed for AM samples (~170 emu/g) are higher than those of the RQ samples (~160 emu/g), these variations are close to the error bar and a dependency cannot be clearly established.

The decrease of the coercivity with the annealing process for the two studied types of samples could be due to a relaxation of the internal stresses of the particles.

However, the different absolute values for RQ and AM samples (both encapsulated powder and pressed bulk disks) is not easy to interpret due to the large number of parameters which define the system: micrometric powder composed of nanocrystalline α-Fe plus other crystalline phases (NbFeB and Fe 2 Zr) in the case of the AM system.

It is worth noting that the trend of H C is not consistent with the evolution of the nanocrystal size in the frame of Herzer's model [START_REF] Herzer | [END_REF]. In fact, samples with the same value of D show very different H C values: e.g. RQ samples after 5, 10 and 15 h of milling and AM samples after 50 and 70 h of milling (see tables 1 and2). On the other hand, after annealing, an increase of D is observed but a decrease of H C is measured.

Therefore, the averaging out of magnetocrystalline anisotropy due to the very fine microstructure predicted by Herzer's model is not the main fact describing the evolution of the coercivity of the powder samples studied in this work and other explanation is required.

The evolution of the size of the powder particles might affect the coercivity of the system. Images of the powders obtained by SEM show that for both RQ and AM samples, the maximum of the particle size distribution saturates at ~25 µm as milling characteristic SEM images are given in Fig. 7, corresponding to RQ sample milled for 10 h and AM sample milled for 50 h. The latter shows a more homogeneous particle size, related to the longer milling time. Therefore, the monotonous evolution of particle size is not sufficient to explain the observed behaviour of coercivity, which is not monotonous (see Fig. 6). Moreover, the differences in coercivity between the two sets of samples can not be explained either by the different particle size of the powder: a higher coercivity is obtained for RQ samples, which have bigger particles.

'[Insert figure 7 around here]'

On the other hand, coercivity might be affected by the magnetoelastic contribution to the anisotropy. In fact, this could explain the differences observed between the two set of samples. AM samples exhibit a lower microstrain than RQ samples, which could yield lower magnetoelastic anisotropy for similar values of magnetostriction. Besides, the presence of several magnetic phases in the AM system could reduce the magnetostriction of the whole system by averaging if the signs of the contributions of the different magnetic phases are opposite: the magnetostriction of Fe 2 Zr is -11 ppm [14] and, although the value of magnetostriction for pure α-Fe is -9 ppm it becomes positive for small Co content [START_REF] O'handley | Modern magnetic materials: principles and applications[END_REF].

On the other hand, the disappearance of intermetallics (rich in Zr, Nb and B) after milling RQ samples implies that the nanocrystal boundaries may be enriched in Zr, Nb and B, non-magnetic atoms insoluble in the α-Fe phase. The effect of a nanocrystal boundary enriched in these atoms, as it might occur in RQ samples, could provoke a more deleterious effect on the magnetic properties through an impingement of the exchange coupling between nanocrystals [START_REF] Hernando | [END_REF] than that produced by the presence of intermetallic inclusions in AM samples through domain wall pinning.

Figure 1 .

 1 Figure 1. The microstructure of the different samples was studied using a Philips PW

  60 degrees. The main contribution corresponds to the (110) diffraction maximum of α-Fe, fitted as a Pseudovoigt function. For a short milling time, tiny crystalline peaks of intermetallic phases appear in this 2θ range (fitted as Lorentzian functions for simplification). In the case of RQ samples, an amorphous halo (simulated as a Gaussian peak) was also necessary to fit the XRD profile. In the case of the AM samples, the Lorentzian contribution to the Pseudovoigt function increases from 0.3 up to 0.6 with the milling time and, after heating up to 973 K, decreases to 0.3, independently of the milling time. This indicates that the relative effect of the small grain size on the broadening of the diffraction maxima increases with the milling time in comparison with the microstrain contribution and decreases with the annealing

  broadening of the (110) diffraction maximum, whereas AM samples do not show significant change.

  for both set of samples. The value of pure α-Fe phase is also shown for comparison.

  For all the samples, the smaller particles are ~5 µm in size. Two

Fe 2 Figure 1 .

 21 Figure captions

Figure 2 .

 2 Figure 2. XRD patterns of powder obtained from a melt-spun ribbon after milling for

Figure 3 .

 3 Figure 3. XRD patterns of powder obtained from an arc-melted alloy after milling for

Figure 4 .

 4 Figure 4. Room temperature XRD patterns of RQ samples annealed at 973 K.

Figure 5 .

 5 Figure 5. Lattice parameter, a, as a function of the milling time for both RQ and AM

Figure 6 .

 6 Figure 6. Coercivity, H C , as a function of the milling time for both RQ and AM samples

Figure 7 .

 7 Figure 7. Secondary electrons SEM images of AM sample milled for 50 h (upper) and

  20 K/min showing the melting and solidifying behaviour of the composition Fe78Co5Nb3Zr3B5Ge5Cu1powder obtained from a melt-spun ribbon after milling for different times. 205x292mm (593 x 593 DPI) Room temperature XRD patterns of RQ samples annealed at 973 K. 205x292mm (593 x 593 DPI) Lattice parameter, a, as a function of the milling time for both RQ and AM samples in the as-milled and annealed state. 295x203mm (593 x 593 DPI) Coercivity, H C , as a function of the milling time for both RQ and AM samples in the as-milled and annealed state. 295x203mm (593 x 593 DPI)

Table 1 .

 1 Microstructural parameters

	As-milled	Heated up to 973 K

Table 2 .

 2 Magnetic parameters

	As-milled	Heated up to 973 K

** Measurements performed ribbon flakes. Not real powder and, therefore, excluded for the annealing study.

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Acknowledgements

This work was supported by the Spanish Government and EU FEDER (Project MAT 2004-04618). J.S. Blázquez acknowledges a research contract from the Regional Government of Andalucía (Spain).


Claire Bergman

Christophe Girardeaux

Carine Perrin-Pellegrino

Patrick Gas

Dominique Chatain

Jean-Marie Dubois

Nicolas Rivier

Wetting of Decagonal Al13Co4 and Cubic AlCo Thin Films by Liquid Pb

Keywords: wetting, intermetallic compounds (user supplied): cubic AlCo, decagonal Al13Co4, Scanning Auger Microprobe wetting, scanning Auger microscopy, intermetallic, cubic AlCo, decagonal Al 13 Co 4

ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

DUBOIS 3 , N. RIVIER 4 1. INTRODUCTION Quasicrystals (QCs) were discovered over 20 years ago [1]. Their bulk atomic structures can be described as a relatively compact quasiperiodic stacking of atoms or clusters of atoms in which typical structural defects develop. QC display bulk physical properties, like low thermal and electrical conductivity that are specific to their structure. Moreover, unusual surface properties have been reported, including low friction coefficients, high contact angles with water and good resistance to oxidation [START_REF] Dubois | New horizons in Quasicrystals: Research and Applications[END_REF][START_REF] Dubois | Quasicrystals. An Introduction to Structure, Physical properties and Applications[END_REF]. Because these anomalous surface properties were investigated in air, an environment where aluminium-rich materials are covered with an aluminium oxide layer, it may be debated wether the observed surface properties are due to the oxide layer, or specific to the quasicrystalline structure.

The first studies of clean surfaces of quasicrystals were published in the early 1990's (for a review see [START_REF] Thiel | [END_REF]). They were focusing on the atomic and electronic structures of the surface. In order to get rid of the oxide layer, several methods have been proposed (cleavage, sputtering, annealing,…) [START_REF] Chevrier | Quasicrystals: Current topics[END_REF]. However they alter the quasicrystallinity of the surface or are not suitable for carrying out wetting experiments.

The present study describes the wetting of Pb, a metal with a low melting point, on the surfaces of two compounds made of the same components, Al and Co, but with different crystal structures. They are the CsCl-type crystal, AlCo, and the decagonal phase Al 13 Co 4 . Special care has been taken to prevent the formation of an aluminium oxide layer on the surface of the Al-Co phases before, and during the wetting experiment. Thus, our method has allowed, for the first time, the study in clean conditions of the wetting of a liquid metal on crystalline and quasicrystalline surfaces.

EXPERIMENTAL

An suitable method for studying the wetting of oxidizable surfaces is the dewetting of thin (~100nm) slabs [START_REF] Wang | [END_REF]. First, the films of Al-Co compounds are prepared (as explained below [7]) in an ultra high vacuum deposition chamber. Then, in the same apparatus and without breaking the vacuum, the lead slab is deposited on the Al-Co substrates. This constitutes the original feature of the process. In addition, the lead layer protects the quasicrystalline surface when the samples are transferred to the scanning Auger microprobe (SAM).

In a previous publication [7], we have shown that a stack of Al/Co bilayers of nanometric thickness forms, upon annealing at about 700K, a single-phase decagonal quasicrystalline film with homogeneous microstructure. The multilayers (300nm) were made by successive electron-beam evaporation of high-purity Al (99.9999 at.%) and Co (99.9999 at.%) on an oxidized Si substrate at room temperature. The base pressure during the evaporation was in the 10 -4 -10 -5 Pa range for both metals. The Al and Co layer thicknesses were monitored during evaporation with a quartz balance. The nominal [Al]/[Co] atomic ratio of the sample corresponds to 13/4 or 1/1, respectively. The transformation into Al 13 Co 4 and AlCo was achieved by an appropriate annealing treatment [7].

After preparation, the Al-Co/Pb sample was transferred in air, from the deposition chamber to a CAMECA Nanoscan 100 scanning Auger microprobe, equipped with a STAIB Instruments analyzer, a 6 keV CAMECA Duoplasmatron ion gun, and a home-built heating unit. Both surface composition analysis and imaging of the particles were undertaken within the SAM. This facility is located at the CRMC-N.

The base pressure in the SAM chamber was 1⋅10 -8 Pa. The surface of the Pb film was first cleaned by ion bombardment, to remove surface contamination gathered during sample transfer. Then the sample was gradually heated, with periodic interruptions during which it was lightly sputtered so as to remove any impurities that may have segregated to the surface. The cleaning process was monitored by Auger electron spectroscopy (AES).

Upon heating above the melting point of Pb (600K), the lead film dewetted the substrate by forming droplets. The surface chemistry of the Al-Co films in-between the droplets was examined by AES. No carbon was detected but oxygen is present in small quantities at the surface of the Al-Co films. Fig. 1 displays an Auger spectrum at high energy, in the derivative mode, where the Auger transition of aluminium appears as metallic. Thus we conclude that the oxygen detected at the surface of the Al-Co compounds is present as adsorbed species at the surface of the Al-Co compounds.

The Pb droplets display in the SAM contact angles lower than 90°. To optimize the wetting measurements, the droplets were frozen and the sample was transferred to a field emission gun scanning electron microscope where pictures of the droplets were obtained with a high resolution. Since there was no significant change of the shape of the droplets after solidification, the measured contact angles were taken as those of the frozen-in Pb droplets.

RESULTS AND DISCUSSION

The figures 2 are general views of the two samples at low magnification. They display Pb droplets uniformly distributed on the samples. The contact circles have diameters of several µm. Figure 3 shows details of typical Pb droplets on which the contact angles are measured. The contact angles, averaged over several droplets, are 42±5° and 49±7° on AlCo and Al 13 Co 4 respectively. This difference is not significant given experimental scatter.

The contact angle of a liquid on a homogeneous surface is given by the Young equation:

cosθ = γ sv -γ sl γ lv (1)
where the γ ij are the interface energy of the three interfaces (l = liquid, s = solid and v = vapor).

It should be emphasized that once the contact angle and the surface energy (surface tension) of the liquid are known, Eq. 1 only gives the difference between the solid/liquid and solid/vapor energies, but not the absolute value of the surface energy of the compound.

To our knowledge there are only estimates [9] but no data on the surface energies of intermetallic compounds. However, we can compare the wetting behaviour of Pb (20K above its melting point) on the Al-Co compounds with the behaviour of Pb on Al 2 O 3 and on pure Al and Co.

The contact angles of Pb on alumina single-crystal and on thick native oxide of aluminium have been reported to be 127±3° and 130±5°, respectively [10]. They are much larger than 90°. Comparison of these contact angles with those measured for the two Al-Co compounds show clearly that the Pb droplets do not lie on aluminium oxide, as already suggested by the Auger spectrum of Fig. 1. On the other hand, contact angles of Pb on aluminium single crystal and on polycrystalline film of Al are close to 30±5° [10,11]. It is only on polycrystalline films of Co that the contact angle of Pb, 50±6° [12], is similar to that on AlCo and Al 13 Co 4 . Thus the wetting behaviour of Pb on both Al-Co compounds is closer to that on a surface of a metal of high melting point.

SUMMARY

The aim of this study was to compare the wetting of a low melting point metal on two Al-Co compounds, one with a CsCl crystalline structure and the other with a quasicrystalline decagonal structure. A challenging point was to design a method where the liquid metal is put into contact with the aluminium-rich compound without need of removing any oxide from the substrate. This was done by preparing films of the compounds and by covering them with a thin Pb slab in a UHV chamber. The thinPb slab was subsequently dewetted to form droplets of micrometer size. The dewetting experiment was performed in a scanning Auger microprobe, in which the surface chemistry can be monitored.

The measured contact angle of liquid Pb (45°) on the Al-Co compounds, when compared to that previously obtained on thick native oxide of aluminium (130°), confirms that our experiments were carried out in clean conditions. Our main new result is that the contact angles observed on both Al-Co compounds are similar, wether their structure is decagonal or cubic. Both compounds are aluminium-rich, but they are not wet like a simple metal such aluminium. Further interpretation of our results requires a better description of the surface of these complex compounds. Acknowledgments: The authors acknowledge the financial support of the Ministère délégué à la Recherche (Action Concertée Incitative 2000, Surfaces, Interfaces et Conception de nouveaux matériaux).

  Wetting of Decagonal Al 13 Co 4 and Cubic AlCo Thin Films by Liquid Pb C. BERGMAN 1 , C. GIRARDEAUX 1 , C. PERRIN-PELLEGRINO 1 , P. GAS 1 , D. CHATAIN 2 , J.M.

FormattedFig. 1 :Fig. 2 :

 12 Figures captionsFig.1 : Auger spectrum in the derivative mode, of the substrate Al 13 Co 4 between Pb droplets taken after fragmentation of the Pb layer. It displays the KLL Auger transition of Al.Fig.2 : Low magnification SEM images of the distribution of the Pb droplets lying on the AlCo and Al 13 Co 4 substrates after freezing.

Fig. 3 :

 3 Fig.3 : High magnification SEM images of Pb droplets on the AlCo and Al 13 Co 4 substrates after freezing. Several of such images were used to measure the contact angles.

  Figure 1

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E Belin-Ferré

M-F Fontaine

J Thirion

S Kameoka

A P Tsai

J M Dubois

Electronic structure of leached Al-Cu-Fe quasicrystals used as catalysts

Keywords: quasicrystals, electronic structure, catalysts user supplied): X-ray emission spectroscopy quasicrystals, catalysts, electronic structure, x-ray emission spectroscopy

We used X-ray emission spectroscopy to investigate the electronic structure of various catalysts obtained by leaching a quasicrystalline Al-Cu-Fe compound in 5%Na 2 CO 3 aqueous solution.

The study of Al states confirmed that upon leaching Al is almost totally oxidised in the catalysts. Fe is not as much oxidised probably because it is embedded in Cu-Al 2 O 3 grains. Cu is only partly oxidised. However, from our data it is not excluded that the catalytic activity could be due to a size effect that prevents the formation of a band as in metallic Cu, and gives to the electronic structure an atomic-like character. New measurements should be done from fresh specimens in order to make sure that no avoid further oxidation interferes with the Cu spectral distributions in the catalysts.

Introduction

Al-Cu-Fe quasicrystalline compounds leached in a 5%Na 2 CO 3 aqueous solution exhibit high performance catalytic properties for steam reforming of methanol [1,2]. Annealing in air at 600°C enhances the catalytic activity, whereas it is decreased by further annealing in H 2 atmosphere. To investigate which is the reason for this behaviour, we studied the electronic structure of such catalysts using X-ray emission spectroscopy (XES) This technique which probes partial electronic distributions densities of states (DOS) around each chemical species in a sample. Here, we have analysed Al 3p, Al 3s,d, Cu 3d and Fe 3d states in various specimens as indicated in the next section For details on XES and the electronic structure of quasicrystals see reference [3].

Experimentals

We have studied the following several samples: (i) sample labelled A is a quasicrystalline Al-Cu-Fe Al 63 Cu 25 Fe 121 powder (grains < 150µm) annealed for 24 hours at 800°C for homogenisation purposes, (ii) sample B is powder A leached in a 5%Na 2 CO 3 aqueous solution and washed to remove all alkali traces, (iii) sample C is powder B, annealed for 4 hours in air at 600°C, (iiii) sample D is powder C, annealed at 600°C for 4 hours in H 2 atmosphere. The Cu Kα x-ray diffraction patterns are dominated by the peaks of the quasicrystalline compound, but also display for samples B, C and D diffraction peaks of Al 2 O 3 , CuO and a β-phase. Cu particles of about 10 nm in size were found to be dispersed at the surface of the catalysts. Pure metals, oxides (Al 2 O 3 and Fe 3 O 4 ) and spinel compounds (CuAl 2 O 4 and CuFe 2 O 4 ) were also analysed for comparison purposes

The XES spectral curves were recorded in vacuum spectrometers fitted with a crystalline slab for Al 3p, Cu 3d and Fe 3d states, respectively, or a grating for Al 3s,d states. The powders stuck onto water cooled holders, were irradiated by electrons of energies allowing to excite the x-ray transitions: valence band (VB) →Al 2p 3/2 , VB →Al 1s, VB →Cu 2p 3/2 and VB → Fe 2p 3/2 levels to probe Al 3sd, Al 3p, Cu 3d and Fe 3d states, respectively. The energy resolution was about 0.4 eV for Al spectra and 0.5 eV for Cu and Fe spectra. The XES technique does not display absolute density of states values. Hence each spectral curve is normalised to its own maximum intensity and it is presented in the x-ray transition energy scale.

The XES spectral curves probe

Results

Figure 1 shows the results for Al 3s-d states in the left side panel and for Al 3p states in the right side panel. In both panels, the two full and dashed curves at the same lower ordinate correspond to samples A and Al 2 O 3 , respectively. From bottom to top, the lines with triangles, diamonds and dots refer to samples B, C and D, respectively. In the right side panel, the upper curve corresponds to spinel CuAl 2 O 4 . The shapes of the Al3s-d curves of B, C and D samples (left side panel) are very similar and depart from the curve of sample A. They look like the curve for Al in Al 2 O 3 . In addition they but display a faint feature around 71 eV in the energy range of a maximum of the curve of sample A. Clearly, samples B, C and D are highly oxidised and show traces of non-oxidised material. For sample B, the peak in the range 67-68 eV is broader than in Al 2 O 3 and bent towards low transition energies due to overlap from oxidised and non-oxidised contributions. The traces of non-oxidised material These are less important for samples C and D than for sample B. Yet, for both C and D specimens, the contributions due to non-oxidised material also broaden and slightly shift the peak in the range 67-68 eV towards low energies as compared to pure alumina. The peak at 64 eV is unchanged in all these samples within the experimental accuracy. In this energy range, the intensity of the curve for sample A, not oxidised, is rapidly decreasing.

The shapes of the Al3p curves for samples B, C and D (right side panel), also significantly depart from the curve of sample A and are similar to the curve that of alumina. Hence, again the samples are highly oxidised. Still, the wing on B, C and D curves around 1558-1560 eV, the energy range where the intensity increases up to of the maximum in curve A, is slightly less hollow than in for Al 2 O 3 . Accordingly, there exists some contribution from non-oxidised material in this energy range. The curves of spinel CuAl 2 O 4 is similar to that of and Al 2 O 3 are analogous. However, the peak at 1539 eV is of higher intensity in the spinel than in oxide; this peak arises from a transition from the VB towards O 2s states hybridised to Al 3p states. The difference in intensity of this peak between the two samples gives an indication that their electronic structures although similar, are not identical. The peak at 1539 eV due to a transition from VB hybridised (O2s -Al3p) states →Al 1s, is of higher intensity in the spinel than in oxide. Hence their electronic structures although similar, are not identical.

At this point, it appears that Thus, Al whatever the specimen, except in powder A, is highly oxidised. In bulk fcc Al, the depths sampled by the excitation beam for investigating Al3s-d and Al3p states respectively, are of the order of 0.3 to 0.5 µm and surface contribution, especially contribution of the oxide, amounts to at most a few per cent of the total signal. For powders, it is not straightforward to quantify precisely the importance of oxidation. Since powder grains represent a large surface contribution to XES data, it makes sense to assume that the oxide contribution dominates the spectra. Actually, it could be estimated from the curves to be of the order of 90% the total signal. In our experimental conditions, the bulk Cu sampled depth is approximately 17 nm. Accordingly, one expects that the whole of Cu grains are probed in the catalysts during XES measurements.

The shape of curve A is reminiscent of that of pure Cu, but narrower due to the alloying effect catalysts and may hidden any other specificity.Thus one may expect that the oxide contribution is significant to the spectra of these catalysts and may hide other specificities.

In contrast, the curve of sample D is wider than for the others. In addition, this curve It shows one marked feature at 929.1 ± 0.2 eV, namely close to the value for CuO. One explanation is that the curve of sample D results from the addition of a CuO contribution to the genuine distribution due to of catalyst D. Actually, the intensities of the CuO peaks in the diffraction pattern of sample D are much lower than in sample C. Actually, the CuO peaks in the diffraction pattern of sample D are of lower intensity than in sample C. Here, we must underline that Note that the catalysts were kept in air. Because they are not reactive in the same way, oxidation upon ageing may have taken place and differ from one sample to the other (the catalytic power of sample D against methanol reforming is slightly lower than that of sample C).

Another change of slope is seen at about 928.5 ± 0.2 eV on the D spectral curve (Fig. 2 At this stage, It is difficult to ascertain whether partial oxidation or size effect are responsible for the shape of the spectral curves of the catalytic samples. We ignore whether the nanometric Cu particles have an "atomic-like" behaviour that could be responsible for their catalytic properties. New experiments are under way, using fresh catalysts kept in argon atmosphere to avoid ageing in air and getting better insight into their genuine electronic structure Figure 3 displays the Fe3d curves in the catalysts, powder A, pure bcc Fe, and also oxide Fe 3 O 4 and spinel CuFe 2 O 4 . Many more scans than for sample A were required by all catalysts to achieve equivalent background-to-signel ratio. For all catalysts, to obtain an equivalent background-tosignal ratio, the measurements required many more scans than for sample A. To interpret this observation, we suggest that Fe atoms are embedded inside Cu-Al 2 O 3 grains. This is supported by the fact that the x-ray diffraction patterns show no contribution due to Fe.

Table 2 collects the data corresponding to these measurements the energy values of the maximum of the Lα band (VB→Fe2p 3/2 ), the FWHM, the ratio between Lα and Lβ peaks (this transition Lβ β β β corresponds to VB→Fe2p 1/2 ) and the energy distance between the two Lα and Lβ sub-bands. The spectral curves of catalysts D and C display a marked feature at the energy of the maximum in the oxide Fe 3 O 4 less pronounced for sample C. Sample C also exhibits such a shoulder but less pronounced. As a result, there is a slight Accordingly, there is a faint broadening of the curves with respect to sample B as seen from the FWHM values. Note that oxidation may occur upon ageing of the samples during the thermal treatments of the samples or, since there is no visible Fe oxide contribution in the x-ray diffraction pattern of the catalysts. , it may depend.

Conclusion

Our XES study of Al states has confirmed that Al is oxidised to at least about 90% in the catalysts. Fe is not as much oxidised, probably because it is embedded in Cu-Al 2 O 3 grains. For Cu, it cannot be excluded from our data that the catalytic activity of samples B, C and D arises from a size effect that prevents the formation of an electron band and gives an "atomic-like" character to the electronic structure. To ascertain this interpretation, measurements must be repeated with catalytic samples kept in an inert atmosphere in order to make sure that no further oxidation modifies their Cu spectral distributions. to avoid further oxidation modifies their spectral distributions.

  Al3s-d, Al3p, Cu3d and Fe3d states, respectively owing to the transitions: valence band (VB) → Al2p 3/2 , VB → Al1s, VB → Cu2p 3/2 and VB → Fe2p 3/2 levels. The energy resolution was about 0.4 eV for Al spectra and 0.5 eV for Cu and Fe spectra and the experimental accuracy between 0.1-0.2 eV. As the XES technique does not provide absolute DOS values, each spectral curve is normalised to its own maximum intensity in the x-ray transition energy scale. Pure metals, Al 2 O 3 , Fe 3 O 4 and spinel compounds (CuAl 2 O 4 and CuFe 2 O 4 ) were also analysed for comparison purposes.

Figure 1

 1 Figure 1 shows the curves for Al3s-d and Al3p states.

Figure 2

 2 Figure 2 displays the Cu 3d distribution curves in both panels. In the left side panel, the bottom curves refer to sample A (full line) and pure Cu (dashed line); from bottom to top, one finds the curves of samples B (triangles), C (diamonds) and D (dots). The vertical dashed line marks the energy position of the maximum for curve A. The right side panel shows the curve of sample D and that of spinel CuFe 2 O 4. .Figure 2 displays the Cu3d distribution curves in the various samples. Within experimental accuracy, the curves of spinels CuFe 2 O 4 and CuAl 2 O 4 are identical.

  [3]. The curves of the catalysts exhibit faint satellites at the same energies as in Cu 2 O [4] for (sample B) or and as in both Cu 2 O and CuO [4] for (sample C). The shape of curve D is more like that of CuO [4] (Insert in Fig. 2 right panel) with a clear satellite around 935 eV. In the spinels no satellite is observed (fig.3 right side panel) which stresses a different electronic structure than in oxides. Because of the satellites, for comparing the full widths at half maximum intensity of the curves (FWHM), one has to consider the low energy part of the spectral distribution and its symmetric with respect to the position of the maximum. The corresponding values are given in Values are in Table 1 together with the energies of the maxima of the curves and the satellites. The data for CuO and Cu 2 O are obtained from reference 3. The data of reference 4 are obtained from the same experimental setup.

  The values for samples A, B, C and D differ from the reference compounds: whereas for compound A the position is quite close to that of Fe in spinel CuFe 2 O 4 , this is no longer true for the other samples. The value of the Lα α α α/Lβ β β β ratio is completely different in compound A from the values for pure Fe, spinel or oxide. The electronic structure of the catalysts differs from that in bcc Fe or Fe 3 O 4 , the Fe oxidation number tends towards that in CuFe 2 O 4 .

Figure1.Figure 2 .

 2 Figure captions

Figure 3 .

 3 Figure 3. Fe 3d distribution curves in bcc Fe (dashed lower curve), Fe 3 O 4 (thin lower line) and spinel CuFe 2 O 4 (thick lower line). At upper coordinates, from bottom to top, are shown the same spectra in samples A (no symbol), B (triangles), C (diamonds) and D (dots). From left to right, the vertical lines mark the maximum of the Fe Lα distribution (see text) in the pure metal, oxide and spinel (dashed line).

Table 1 Sample

 1 Within our experimental precision, except for fcc Cu and sample D, the FWHM of the curves are the same. and the maxima of the curves of the three catalysts are found at intermediate energies between the ones for the two oxides. They differ for the spinels. The

		Max Cu 3d	Energy of Satellites	FWHM
		± 0.15 eV	± 0.2 eV	± 0.1 eV
				± 0.2 eV
	Cu	930.10	932.8 935.7	3.0
	CuO	929.30	931.8 934.7	2.0
	Cu 2 O	929.70	931.2	2.0
	CuAl 2 O 4 (spinel)	930.40		2.0
	CuFe 2 O 4 (spinel)	930.40		2.0
	Quasicrystal : sample A	929.50	931.5 934.4	2.0
	Leached : sample B	929.50	931.2 934.4	2.0
	Leached+annealed : sample C	929.60	931.2 934.6	2.0
	Leached+annealed+H 2 : sample D	929.50	928.8 931.6 934.6	2.6

maxima of the curves of the three catalysts differ from the spinels and are at energies intermediate between those for the two oxides

	F o r
	P
	e
	e r
	R
	e
	v i e
	w
	O n l
	y

. From the x-ray diffraction pattern show CuO peaks, should contribute to the spectra especially in sample C. The FWHM are the same in catalysts B and C and in oxides. Because the maxima are at the same energy value one may expect that the oxide contribution dominates the spectra of these

  right panel. Hence, this curve shows 3 features. It is known that in Ni nano-sized metallic particles, the 3d ban, d is no longer metallic-like, but shows 5 features which are ascribed to the 5 d individual orbitals[5]. A similar behaviour may occur in the Cu catalysts that are less than 10 nm in size but this effect may be hidden because of oxidation.

Table 2

 2 The values for samples A, B, C and D differ from the reference compounds, the maxima are shifted opposite to the values of the pure metal. Whereas for compound A the position is quite close to that of Fe in CuFe 2 O 4 spinel, this is no longer true for the other samples. The value of the Lα/Lβ ratio is completely different in compound A from the ones for pure Fe, spinel or oxide. The trend shown by the catalysts is to tend towards the values for the spinel. In other words, the electronic structure in the catalysts is completely different from that in bcc Fe or Fe 3 O 4 , but could be similar to that in CuFe 2 O 4 .

	Sample	Max Lα	FWHM	Lα/Lβ ratio	Lβ-Lα
		± 0.15 eV	± 0.1 eV	± 0.2	± 0.2 eV
			± 0.2 eV
	Fe	704.95	5.1	2.1	12.4
	Quasicrystal : sample A	706.54	4.4	3.3	13.0
	Leached : sample B	706.78	5.0	2.5	11.4
	Leached+annealed : sample C	707.56	5.4	2.3	11.3
	Leached+annealed+H 2 : sample D	707.32	5.4	2.2	11.6
	Fe 3 O 4	705.40	5.2	2.1	12.0
	CuFe 2 O 4 (spinel)	706.50	5.4	1.9	11.4

Table captions Table 1 .

 captions1 FWHM and energies of the features of Cu 3d spectral curves

Table 2 .

 2 Energies values of the features of the Fe 3d spectral curves

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