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The oxidation at 500 °C of AlCuFe quasicrystalline powders: A X-ray diffraction study
Journal of Non-Crystalline Solids 351 (2005) 1630–1638 www.elsevier.com/locate/jnoncrysol
The oxidation at 500 C of AlCuFe quasicrystalline powders: A X-ray diffraction study P. Weisbecker a
a,*
, G. Bonhomme
a,b
, G. Bott a, J.M. Dubois
a
Laboratoire de Science et Ge´nie des Mate´riaux et de Me´tallurgie (UMR7584 CNRS-INPL), Ecole des Mines, Parc de Saurupt, 54042 Nancy, France b Coatings Solutions, Saint-Gobain NRDC, Northborough, MA 01532, United States Received 16 September 2004; received in revised form 6 April 2005
Abstract The oxidation behavior at 500 C of AlCuFe icosahedral quasicrystalline powders was investigated by means of X-ray diffraction (XRD), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Powders have been characterized at room temperature before and after oxidation. Oxidation leads to a long-range diffusion of aluminum atoms, which form an amorphous oxide layer on top of the grains. When the loss of aluminum inside the grains is too high, the icosahedral phase transforms in cubic b-phase (CsCl type). This transformation no longer occurs when the grain size is higher than a critical value. 2005 Elsevier B.V. All rights reserved. PACS: 61.44.B; 81.65.M; 61.10.N
1. Introduction Most potential applications of quasicrystals are in the field of coatings. The processing route to prepare these coatings includes the preparation of atomized powders whose oxidation state can drastically change the properties of the industrial product. Moreover, powders are very sensitive to oxidation due to their high specific area, which in turn makes them suitable to study the influence of oxidation. The aim of this paper is to bring an answer to the following questions: What is the nature of the oxide layer, what is its growth kinetics?
* Corresponding author. Address: LCTS, Domaine Universitaire, 3 Allee de La Boetie, 33600 Pessac, France. Tel.: +33 5 5684 4718; fax: +33 5 5684 1225. E-mail address: [email protected] (P. Weisbecker).
0022-3093/$ - see front matter 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2005.04.053
Is the quasicrystalline phase destabilized by oxidation and does it transform into an approximant phase? Does oxygen diffuse into the icosahedral lattice? Several studies were carried out by various groups in the AlCuFe ternary system [1–4]. For temperatures under 300 C, it has been shown that a thin amorphous oxide layer grows at the surface. For higher temperatures, 800 C for instance, Wehner et al. [5] have shown that the early oxidation stage is determined first by the growth of c-Al2O3, then needle-shaped h-Al2O3 appears, last a-Al2O3alumina nucleates at the metal–oxide interface and randomly grows through the h-Al2O3 layer. Copper seems to trigger the transformation from h-Al2O3 to a-Al2O3 that usually takes place at higher temperatures. At intermediate temperatures, for instance at 500 C, Yamasaki and Tsai [6] have shown that a phase transformation from the i-phase to a CsCl-type cubic Al(Cu,Fe) phase (b 0 -phase) is triggered by oxidation. However, information concerning the growth and the
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nature of the oxide layer are missing, especially about the kinetics of the phase transformations and the diffusion of oxygen in the lattice. As a consequence we have studied oxidation under air at 500 C of several samples of i-AlCuFe powders. Evolution of structural and micro-structural parameters was mainly followed by in-lab X-ray diffraction, whereas selected samples were characterized by SEM and TEM.
2. Experimental 2.1. Sample preparation and oxidation conditions Pure elements, Al (99.999%), Fe (99.97%), Cu (99.99%) are melted by induction in order to obtain an ingot. This ingot is then ball milled using steel jars and balls. In order to avoid contamination from the iron coming from the balls and jar, ball-milling cycles are limited to 2 min. Between each cycle, an amount of powder with particle size falling in the range 25–53 lm is removed from the miller. Roughly 20 g of powder is then sintered under uniaxial load. In order to avoid the oxidation of the alloy, the powder is first heated at 300 C under primary vacuum (5 · 103 mbar) then the full load (1000 daN) is applied on the sample under helium atmosphere. Next, the sample is maintained during 1 h at the sintering temperature (800 C). After sintering, the samples are polished in order to remove the graphite layer from the surface, they are annealed at various temperatures and quenched in air or water. After heat treatment, samples are once again polished to remove the oxide layer and finally crushed to obtain powders suitable for the oxidation experiment. Different particle size distributions are obtained by sieving. Table 1 shows the preparation and oxidation conditions of the samples used in this study. Various compositions have been chosen around the ÔidealÕ composition Al62.5Cu25.3Fe12.2 (at.%). Indeed, XRD and TEM observations have shown that this composition leads to a pure icosahedral phase of high structural quality. Aluminum content was slightly changed in
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the other samples so as to stay in the stability domain of the quasicrystalline phase, whereas iron and copper compositions were adjusted to keep a Cu/Fe ratio close to 2.07. SEM images (Fig. 1) show that the grains have a similar irregular shape whatever the particle size. Previous to the oxidation experiments, each batch was characterized by XRD and TEM. The only observed phases were the pure icosahedral quasicrystal for compositions Al63Cu25Fe12 and Al62.5Cu25.3Fe12.2 and a mixture of icosahedral quasicrystal and of b-phase ˚ ) for the composition Al62(cubic, CsCl type, a = 2.9 A Cu25.6Fe12.4. In this case, the amount of b-phase was estimated to be lower than 5 wt%. The structural perfection of the quasicrystals can be estimated thanks to XRD peak broadening. Hence sample Al63Cu25Fe12 exhibits broad peaks and TEM has shown that this a modulated quasicrystal; quenching allows to obtain narrower peaks in the XRD diagram but the modulation is preserved. Other compositions exhibit narrow peaks by XRD and the structural quality of the quasicrystalline phase is good as observed by TEM
Fig. 1. SEM image from grains (white) belonging to the 25–53 lm particle size granulometry after crushing in an agate mortar.
Table 1 Compositions, samples preparation details and oxidation conditions Nominal composition (at.%)
Heat treatment
Phases identified before oxidation
Particle size (lm)
Oxidation conditions
Al63Cu25Fe12
800 C water quenched
Modulated icosahedral quasicrystal
<25
730 C water quenched
Modulated icosahedral quasicrystal
25–53 <25 25–53
500 C 500 C 500 C 500 C 500 C 500 C
Al62.5Cu25.3Fe12.2
730 C air quenched No heat treatment
Icosahedral quasicrystal Icosahedral quasicrystal
<25 <25
500 C in air 500 C in air
Al62Cu25.6Fe12.4
730 C air quenched
Icosahedral quasicrystal and b-phase.
<25
500 C in air
at at in in in in
1 · 104 mbar 1 · 106 mbar air air air air
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imaging and evidenced by the sharpness of the diffraction spots. Most samples were oxidized at 500 C in air using a conventional furnace. Powders were placed in an alumina container. In order to check the reproductibility of the experiment, it was done twice at two months interval for the sample having composition Al63Cu25Fe12. Moreover, some experiments were carried out under primary vacuum (1 · 104 mbar) and second– ary vacuum (1 · 106 mbar) in order to analyze the influence of the annealing conditions with and without an oxidatizing atmosphere. All samples were characterized at room temperature after cooling down from 500 C.
peak asymmetry and the amount of b-phase were then deciphered from the X-ray pattern. In order to determine the six-dimensional lattice parameter a6d we have used the indexing scheme proposed by Cahn et al. [8]. The AlCuFe icosahedral phase can be ascribed a Face Centered Cubic lattice in a sixdimensional real space with lattice parameter a6d. A simple relationship exists between inter-planar distances, lattice parameter a6d and diffusion vector qk: pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi sin h 2p N þ sM ˚ 1 Þ; pffiffiffiffiffiffiffiffiffiffiffiffiffi ðA ¼ jqk j ¼ 4p k a6d 2 þ 2s
A Siemens D500 diffractometer (Co anode, kKa1 = ˚ ) was used to characterize the powders. It is 1.78897 A fitted with a primary bent monochromator and a linear detector. The XRD patterns were recorded for 2h ranging from 25 to 120. The step size was equal to 0.032 2h, which corresponds to a channel of the detector, and the counting time was fixed to 1 h. In order to determine accurately the six-dimensional lattice parameter of ˚ ) was the icosahedral phase silicon powder (a = 5.4309 A used as an internal standard to obtain the experimental errors. Fig. 2 shows a powder X-ray diagram for the composition Al62.5Cu25.3Fe12.2 before oxidation. This diffraction pattern is representative of an icosahedral material of high lattice perfection. Diffraction peaks due to the internal silicon standard are also indicated. The X-ray diffraction patterns were fitted with an asymmetric Pearson VII profile shape function using the computer program Winfit [7]. The six-dimensional lattice parameter, FWHM (full width at half maximum),
where N and p Mffiffiffi are the indices of the diffracting planes and s ¼ ð1 þ 5Þ=2 is the golden number. The lattice parameterpffiffiffiffiffiffiffiffiffi is deduced from XRD peak ffi N þsM positions by plotting 2ppffiffiffiffiffiffiffi versus q k. One therefore 2þ2s obtains a straight line whose slope equals 1/a6d. The a6d parameter is a measure of the six-dimensional unit cell edge length. However, the three-dimensional quasiperiodic structure being obtained by a cut of 6D space, changes of a6d have the same meaning than in conventional crystallography. Thus, we can obtain information about the unit cell dimension and the atomic density, therefore about the formation of vacant sites, about structural changes due to enrichment or loss of an element (chemical change) and oxygen diffusion in the lattice. The relative changes in the FWHM were measured for peaks (70/113) and (72/116) as a function of oxidation time. As it is well known, peak broadening is linked with the occurrence of micro-strains and with crystallite sizes. However, for quasicrystals a relationship exists in the presence of a phason strain field between peak broadening and diffusion vector in perpendicular space q? that allows to quantify the strength of the phason field [9–12]. Fig. 3 shows the two peaks considered to obtain reliable data along the lines suggested.
Fig. 2. X-ray pattern for Al62.5Cu25.3Fe12.2 powder. Silicon peaks are pointed out (vertical bars) and indices for the main peaks of the icosahedral phase are given.
Fig. 3. Peaks (70/113) and (72/116) for a non-oxidized sample. The FWHM is close to the instrumental broadening.
2.2. X-ray diffraction characterization
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– and its lattice parameter for composition Al50Cu34Fe16 ˚ . This phase often coexists with the (at.%) is ab = 2.942 A i-AlCuFe phase because it forms during cooling from the liquid and participates to the peritectic reaction which causes the formation of the icosahedral phase. A ratio named F(b) was used as a marker to follow the growth of the b-phase. To define F(b), we used the intensities of the b-phase (1 1 0) peak and of the icosahedral (18/29) peak: F ðbÞ ¼ 100
Fig. 4. Icosahedral phase (72/116) peak used to quantify the peak asymmetry and (5 1 1) silicon peak.
This ratio does not give directly the b-phase volume fraction. However it can be calculated using the following formula: n¼
An asymmetry parameter a was defined as the difference between the half width at half maximum left (HWHMleft) and the half width at half maximum right (HWHMright): a = HWHMleft HWHMright (2h). This parameter has been calculated for peak (72/116). It is mainly due to a lattice parameter gradient, which can be inter-particles or intra-particles. If this gradient is only intra-particles, it can be due to oxygen or aluminum diffusion toward the grain surface. It must be emphasized that the penetration depth of X-rays at Co Ka wavelength in this AlCuFe alloy is around 20 lm. Fig. 4 shows the asymmetry of peak (72/116) of the icosahedral phase in an oxidized sample. One peak of the silicon standard is also shown. In one sample the b-phase (see Fig. 5 for indexation) is present previous to oxidation. It is a solid solution of composition Al50x(Cu,Fe)50+x which occupies a very broad field in the ternary system AlCuFe since nearly all iron can be substituted by copper. The b-phase structure belongs to the space group Pm 3m – CsCl structure
Ið1 1 0Þ . Ið18=29Þ
K Ið1 1 0Þ ; K Ið1 1 0Þ þ Ið18=29Þ
where n is the b-phase volume fraction and K is a constant determined by calibration on a sample containing 90%w of icosahedral phase and 10%w of b-phase powder (K 0.06).
3. Results The aim of this work was to study powder oxidation. However at 500 C, other mechanisms that would influence the evolution of various parameters can take place, namely phase transformation if equilibrium was not reached during cooling, crystallite growth and coalescence, etc. In order to study the effects of oxidation, we needed to make sure which changes are only due to oxidation. With this goal in mind, several samples were annealed at 500 C in a furnace under high vacuum (1 · 106 mbar). The evolution of lattice parameters in this condition is highly slowed down since the partial oxygen pressure is low. For all batches annealed in such conditions (composition Al63Cu25Fe12 at.%, annealed at 800 C, water quenched, particle size lower than 25 lm), a 178 h annealing causes a change in parameters, for instance a6d, similar to 2 h oxidation in air. After 178 h ˚ whereas annealing under vacuum, a6d = 6.3159 A ˚ a6d = 6.3156 A after 2 h in air. Thus it appears obvious that the parameter changes are mainly due to grain oxidation. 3.1. Evolution of the six-dimensional lattice parameter a6d
Fig. 5. Peaks positions and Miller indices for the main reflections of the b-phase: (1 0 0), (1 1 0), (2 1 1).
3.1.1. Composition effect Numerous papers dealing with quasicrystals give values for the six-dimensional lattice parameter but due to the various processing routes and characterization methods used, they can hardly be compared. However Quiquandon et al. [13] carried out an extensive study on
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the influence of the composition on the lattice parameter. They showed that the six-dimensional lattice parameter is proportional to the aluminum atomic content at constant electron/atom ratio (e/a). Using their experimental results we can infer a relationship between a6d, the Al at.% and the (e/a) ratio: a6d ¼ 0.3698 ðAl at:%Þ 0.1255 ðe=aÞ þ 6.32156. ð1Þ In our study the three selected compositions were chosen in order to keep the same ratio between copper and iron contents. Its is interesting to see that the lattice parameter is inversely proportional to the aluminum content. This behavior could be expected from Eq. (1) as can be seen in Fig. 6, however the slope is higher than expected and the difference increases with the aluminum content. This discrepancy could be due to the difference between the nominal composition of our samples and the real composition as well as to the presence of impurities as oxygen or carbon introduced during the sintering process. During oxidation the trend is the same whatever the composition, the initial heat treatment and the particle size. In all cases, the lattice parameter increases with oxidation time. This trend is shown in Fig. 7 for the three compositions, for particle sizes lower than 25 lm and for initially quenched samples. Using a linear time scale, the increasing lattice parameter in six-dimensional space seems to reach saturation beyond a few hundred hours. However, a6d evolution is controlled by atomic diffusion, which is governed by nonlinear equations. In this way, using a logarithmic time scale allows us to correct this effect. It appears that all series have a similar evolution: a rather linear increase of a6d with time in the beginning whereas it is only for long oxidation times
Fig. 7. a6d evolution with oxidation time at 500 C. The particle size is lower than 25 lm for the three series. Linear (a) and logarithmic time scale (b). Lines are only given to guide the eye. Error bars are typically of the same size than symbols.
that saturation comes in (higher than 1000 h of oxidation time). A slight decrease seems also to be observed for the largest oxidation times. 3.2. Particle size effect In order to understand the effect of particle sizes, the time dependence of the events is summarized in Fig. 8 in terms of lattice parameter for two different particle sizes. To check the reproducibility of the experiment, it was carried out twice for the Al63Cu25Fe12 batch. The lattice parameter increases more quickly for the lower particle sizes whereas the final lattice parameter is larger, which confirms that the process is linked with the specific area, because the effects of oxidation are more pronounced near the surface and are all the more amplified when the ratio area to volume is the largest. 3.3. Peaks width and asymmetry evolution
Fig. 6. Variation of the 6D lattice parameter a6d with the alloys aluminum content, before oxidation and comparison with Quiquandon [12] results. The continuous lines gives a6d using the equation: a6d = 0.3698 · (Al at.%) 0.1255 · (e/a) + 6.32156, derived from Quiquandon results.
The peaks full width at half maximum (FWHM) and asymmetry behave in a similar way as can be seen in Fig. 9. They both increase during the first step of oxidation, then reach a maximum and afterwards decrease to val-
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Fig. 8. Grain size effect on a6d evolution during oxidation at 500 C for powder Al63Cu25Fe12 – the second series of measurements was made several months after the first one.
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Fig. 10. Increase of the b-phase amount with oxidation time at 500 C for the composition Al63Cu25Fe12, particle size bellow 25 lm. Evaluation of the annealing temperature effect and of the reproducibility. The evolution of the (1 1 0) b-phase peak can be seen at the bottom of the graphic.
tion, whatever the composition may be, this phase systematically appears in amounts that increase with oxidation time as can been seen in Fig. 10 for several series. The richer in aluminum the alloy is, the later the b-phase appears and the lower the b-phase amount at saturation is. However, the various compositions follow all the same trend that is characterized by an increase at the beginning of oxidation and a saturation after an inflection point for longer oxidation times. The influence of the particle size on the b-phase amount can be assessed in Fig. 11. This variation is amplified when the particle size is the lowest, as it was already pointed out for the a6d parameter of the icosahedral phase. It must be emphasized that the lattice parameter of the b-phase does not change significantly with oxidation ˚ . FWHM and time and lies between 2.94 and 2.945 A asymmetry parameters exhibit similar behavior. From this result, it is straightforward to deduce that the bphase does not take part in the oxidation process, but is just an oxidation product. 3.5. Composition of the oxide layer
Fig. 9. FWHM (a) and asymmetry parameter a (b) evolution with oxidation time at 500 C for two particle sizes.
ues close to their initial values. Changes are more drastic for the smallest particle sizes. 3.4. Analysis of the b-phase Before oxidation, the b-phase is present in only one alloy, the poorest in aluminum. However during oxida-
The powders show a steel gray color that becomes darker and bluish with oxidation time. It is reasonable to assume that it is essentially aluminum, which oxidizes and iron in a lesser extent. Even for long oxidation times – three months at 500 C – no other crystalline phase than the b-phase and the icosahedral phase was detected. Particularly, no crystallized oxide appeared on the surface of the grains. From X-ray diffraction results it is not possible to conclude that the oxide layer is completely amorphous. Indeed the phase fraction could be too small to give rise to measurable peaks by XRD and diffraction peaks could be broadened by domain size effects. However a TEM study, whose results will be presented elsewhere, carried out on bulk samples, under the same oxidation conditions, showed that a
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Fig. 11. F(b) parameter variation with oxidation time at 500 C (a) for the two particle sizes for composition Al63Cu25Fe12 and (b) for various compositions (particle size bellow 25 lm). A logarithmic time scale is used.
homogenous and amorphous oxide layer appears on the sample surface during oxidation. This result is not clearly understood now because usually in aluminum alloys, amorphous aluminum oxide, which appears in the first step of the oxidation, quickly crystallizes for temperatures above 400–450 C [14].
asymmetry of the peaks can be due to a lattice parameter gradient and their values give information about the gradient intensity. The a6d parameter increases when the aluminum content decreases in the alloy. The same is observed when the b-phase appears and its amount increases. So, the b-phase appears when the loss of aluminum in the icosahedral phase is too significant, which is consistent with the phase diagram at 500 C. This aluminum loss is not uniform and even for the composition the richest in aluminum, it is likely that some zones quickly reach the 1 at.% loss in aluminum necessary to trigger the phase transformation. It is likely that the surface oxide layer is mainly composed of aluminum oxide with a small amount of iron in solution and no copper. The three compositions were chosen so that they stand on a straight line in the ternary phase diagram. During the oxidation process, the composition of the icosahedral phase probably follows this line, which was expected if only aluminum contributes the oxide layer. The plot of a6d with the b-phase ratio confirms this assumption since it can be seen in Fig. 12 that all compositions follow the same trend. This trend is linked with the [Cu]/[Fe] ratio and with the aluminum reservoir in the sample, i.e. with its specific area. This implies that for a given composition, the stability of the icosahedral phase depends on the quantity of aluminum exposed to contact with oxygen, as characterized by the surface area to volume ratio (S/V) of the particle. This effect is possible if aluminum diffusion is a long distance mechanism, at least over distances larger than 50 lm. If the diffusion was only a surface phenomenon, localized to within a few microns, no difference would be observed between different particle sizes. It must also be underlined that oxidation kinetics are similar whatever the size of the aluminum reservoir is. This kinetics slows down and seems to follow a parabolic evolution. After 1000 h oxidation, the oxidation rate is very low on a linear time scale and the phenomena seem to reach a saturation point.
4. Discussion 4.1. The aluminum reservoir effect In order to understand the evolution of the various parameters, it is interesting to point out the correlation that exists between each other. Whereas the width and the asymmetry of the diffraction peaks show similar evolution, the same is true for the a6d lattice parameter and the amount of b-phase, which by the way seem closely related. Moreover the first two parameters reach their maximum value at the half duration of the two latter parameters evolution, which corresponds to an inflection point. Thus, the increase of the width and of the
Fig. 12. a6d parameter evolution with b-phase fraction (F(b) parameter). Particle size bellow 25 lm.
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Fig. 14. Avrami plot for the sample Al62.5Cu25.3Fe12.2 (at.%), particle size <25 lm, oxidized at 500 C in air. Fig. 13. Plot showing the Al depletion in function of the oxide layer thickness for the sample Al63Cu25Fe12 (at.%) with four different particle sizes: 5, 25, 53 and 100 lm. The model uses spherical powder particles, and a homogeneous volumic depletion. The model also takes into account the atomic density differences (in Al at. nm3) between alumina and the quasicrystalline matrix. The dotted line at 62.2 at.% Al corresponds to the limit of the stability range of the quasicrystalline phase.
The Al depletion has been computed in function of the oxide layer thickness for different particle sizes (Fig. 13). Since the real S/V ratio could not be measured for the different samples, the model assumes the powder particles have a spherical shape, which is a rough approximation and leads to underestimate the Al depletion. Nonetheless, the diagram shows that for particle sizes larger than 100 lm it is unlikely that the icosahedral to b-phase transformation would occur with the oxidation conditions studied here.
Three processes pictured in Fig. 15 can be expected for the formation and growth mechanisms of the oxide. Each grain is embedded in a native oxide layer which forms during crushing. At 500 C, this layer grows following a parabolic kinetics because it implies aluminum diffusion toward the oxide/atmosphere interface. The diffusion coefficient being larger at the grain boundaries than in the bulk, it is likely than the aluminum depletion is more important at the boundary between two icosahedral grains and the oxide. Nucleation of the b-phase could possibly occur at such sites. The platelets could then grow from this interface towards the core of the icosahedral phase crystallites. Another process could lead to the formation of plates at metal/oxide interface by transformation of the icosahedral phase grains being
4.2. Kinetics of the b-phase transformation: an Avrami-law analysis Usually, an isothermal transformation can be analysed using the Avrami equation [15]: n ¼ 1 expðktn Þ; where n is the transformed volume fraction, k is the crystallization rate and the value of n is governed by the transformation mechanism. Christian [15] has given the various transformation mechanisms and corresponding n exponents. Thus, n can be determined using the following equation: 1 log log ¼ n logðtÞ þ logðkÞ. 1n The curve, plotted in Fig. 14 is obtained for the Al63Cu25Fe12 composition. The slope is roughly equal to 1/2, which indicates that the transformation is diffusion controlled, in accordance with the growth of broad platelets at the oxide/metal interface or/and at the grain boundaries where the aluminum depletion is the most significant.
Fig. 15. Schematic model for the three possible oxidation processes described in this paper. Powder grains are shown before, during the first oxidation steps and after a long oxidation time. (a) All crystallites of the grain transform into b-phase. (b) Only the crystallites near the surface of the grain undergo a phase transformation. (c) The transformation occurs at the grain boundaries and the b-phase crystallites look like platelets.
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in direct contact with the oxide layer. The last process, which is not incompatible with the two others, consists in the transformation of a whole grain of icosahedral in b-phase by growth at the grain boundaries or at the metal/oxide interface. The first process is the most plausible, but it is possible that the two nucleation processes coexist and that some grains fully transform in b-phase. 5. Conclusions We have shown that the oxidation behavior of i-AlCuFe powders at 500 C is mainly governed by diffusion of aluminum atoms towards the surface of the powder particles. Aluminum oxidizes and forms an amorphous oxide layer on the surface of the grains whatever the oxidation time. When the aluminum loss is too large, the icosahedral phase transforms into b-phase, in accordance with the phase diagram at 500 C. Another composition of the quasicrystal or another oxidation temperature would lead to a different phase transformation in accordance with the same phase diagram. Aluminum diffusion is a long distance process, which means that for a particle size larger than a critical particle size, appearance of the b-phase, due to a too important depletion in aluminum, can be avoided. This phenomenon was indeed observed in bulk samples, which have a lower specific area and contain a more important aluminum reservoir.
References [1] S.S. Kang, J.M. Dubois, J. Mater. Res. 10 (5) (1995) 1071. [2] D.J. Sordelet, M.J. Kramer, I.E. Anderson, M.F. Besser, in: C. Janot, R. Mosseri (Eds.), Proceedings of the 5th International Conference on Quasicrystals, World Scientific, Singapore, 1995, p. 778. [3] D.J. Sordelet, L.A. Guderman, M.F. Besser, A.B. Akinc, in: New Horizons in Quasicrystals: Research and Applications Conference Proceedings, Singapore, 1997, p. 296. [4] P. Weisbecker, G. Bonhomme, A. Cael, L. Zhang, J.M. Dubois, Mater. Res. Soc. Symp. Proc. K9 (2) (2001) 1. [5] B.I. Wehner, U. Ko¨ster, A. Ru¨diger, C. Pieper, D.J. Sordelet, Mater. Sci. Eng. A 294–296 (2000) 830. [6] M. Yamasaki, A.P. Tsai, J. Alloys Compd. 342 (2002) 473. [7] S. Kru¨mm, in: XIIIth, Conference on Clay Mineralogy and Petrology, Praha, 1994. [8] J.W. Cahn, D. Shechtman, D. Gratias, J. Mater. Res. 1 (1986) 13. [9] M. Jono, Y. Matsuo, Y. Ishii, Mater. Sci. Eng. A 294–296 (2000) 680. [10] M. de Boissieu, H. Takakura, M. Bletry, J.Q. Guo, A.P. Tsai, J. Alloys Compd. 342 (1–2) (2002) 265. [11] A. Le´toublon, I.R. Fisher, T.J. Sato, M. de Boissieu, M. Boudard, S. Agliozzo, L. Mancini, J. Gastaldi, P.C. Canfield, A.I. Goldman, A.P. Tsai, Mater. Sci. Eng. A 294–296 (2000) 127. [12] A.P. Tsai, A. Inoue, T. Masumoto, Prog. Cryst. Growth Char. 34 (1997) 221. [13] M. Quiquandon, A. Quivy, J. Devaud, F. Faudot, S. Lefebvre, M. Bessie`re, Y. Calvayrac, J. Phys.: Condens. Matter 8 (1996) 2487. [14] K. Wefers, C. Misra, Oxides and hydroxides of aluminum, Alcoa Technical Paper no. 19, Alcoa Laboratories, 1987. [15] J.W. Christian, The Theory of Transformations in Metals and Alloys, Pergamon, London, 1965.
The oxidation at 500 C of AlCuFe quasicrystalline powders: A X-ray diffraction study P. Weisbecker a
a,*
, G. Bonhomme
a,b
, G. Bott a, J.M. Dubois
a
Laboratoire de Science et Ge´nie des Mate´riaux et de Me´tallurgie (UMR7584 CNRS-INPL), Ecole des Mines, Parc de Saurupt, 54042 Nancy, France b Coatings Solutions, Saint-Gobain NRDC, Northborough, MA 01532, United States Received 16 September 2004; received in revised form 6 April 2005
Abstract The oxidation behavior at 500 C of AlCuFe icosahedral quasicrystalline powders was investigated by means of X-ray diffraction (XRD), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Powders have been characterized at room temperature before and after oxidation. Oxidation leads to a long-range diffusion of aluminum atoms, which form an amorphous oxide layer on top of the grains. When the loss of aluminum inside the grains is too high, the icosahedral phase transforms in cubic b-phase (CsCl type). This transformation no longer occurs when the grain size is higher than a critical value. 2005 Elsevier B.V. All rights reserved. PACS: 61.44.B; 81.65.M; 61.10.N
1. Introduction Most potential applications of quasicrystals are in the field of coatings. The processing route to prepare these coatings includes the preparation of atomized powders whose oxidation state can drastically change the properties of the industrial product. Moreover, powders are very sensitive to oxidation due to their high specific area, which in turn makes them suitable to study the influence of oxidation. The aim of this paper is to bring an answer to the following questions: What is the nature of the oxide layer, what is its growth kinetics?
* Corresponding author. Address: LCTS, Domaine Universitaire, 3 Allee de La Boetie, 33600 Pessac, France. Tel.: +33 5 5684 4718; fax: +33 5 5684 1225. E-mail address: [email protected] (P. Weisbecker).
0022-3093/$ - see front matter 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2005.04.053
Is the quasicrystalline phase destabilized by oxidation and does it transform into an approximant phase? Does oxygen diffuse into the icosahedral lattice? Several studies were carried out by various groups in the AlCuFe ternary system [1–4]. For temperatures under 300 C, it has been shown that a thin amorphous oxide layer grows at the surface. For higher temperatures, 800 C for instance, Wehner et al. [5] have shown that the early oxidation stage is determined first by the growth of c-Al2O3, then needle-shaped h-Al2O3 appears, last a-Al2O3alumina nucleates at the metal–oxide interface and randomly grows through the h-Al2O3 layer. Copper seems to trigger the transformation from h-Al2O3 to a-Al2O3 that usually takes place at higher temperatures. At intermediate temperatures, for instance at 500 C, Yamasaki and Tsai [6] have shown that a phase transformation from the i-phase to a CsCl-type cubic Al(Cu,Fe) phase (b 0 -phase) is triggered by oxidation. However, information concerning the growth and the
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nature of the oxide layer are missing, especially about the kinetics of the phase transformations and the diffusion of oxygen in the lattice. As a consequence we have studied oxidation under air at 500 C of several samples of i-AlCuFe powders. Evolution of structural and micro-structural parameters was mainly followed by in-lab X-ray diffraction, whereas selected samples were characterized by SEM and TEM.
2. Experimental 2.1. Sample preparation and oxidation conditions Pure elements, Al (99.999%), Fe (99.97%), Cu (99.99%) are melted by induction in order to obtain an ingot. This ingot is then ball milled using steel jars and balls. In order to avoid contamination from the iron coming from the balls and jar, ball-milling cycles are limited to 2 min. Between each cycle, an amount of powder with particle size falling in the range 25–53 lm is removed from the miller. Roughly 20 g of powder is then sintered under uniaxial load. In order to avoid the oxidation of the alloy, the powder is first heated at 300 C under primary vacuum (5 · 103 mbar) then the full load (1000 daN) is applied on the sample under helium atmosphere. Next, the sample is maintained during 1 h at the sintering temperature (800 C). After sintering, the samples are polished in order to remove the graphite layer from the surface, they are annealed at various temperatures and quenched in air or water. After heat treatment, samples are once again polished to remove the oxide layer and finally crushed to obtain powders suitable for the oxidation experiment. Different particle size distributions are obtained by sieving. Table 1 shows the preparation and oxidation conditions of the samples used in this study. Various compositions have been chosen around the ÔidealÕ composition Al62.5Cu25.3Fe12.2 (at.%). Indeed, XRD and TEM observations have shown that this composition leads to a pure icosahedral phase of high structural quality. Aluminum content was slightly changed in
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the other samples so as to stay in the stability domain of the quasicrystalline phase, whereas iron and copper compositions were adjusted to keep a Cu/Fe ratio close to 2.07. SEM images (Fig. 1) show that the grains have a similar irregular shape whatever the particle size. Previous to the oxidation experiments, each batch was characterized by XRD and TEM. The only observed phases were the pure icosahedral quasicrystal for compositions Al63Cu25Fe12 and Al62.5Cu25.3Fe12.2 and a mixture of icosahedral quasicrystal and of b-phase ˚ ) for the composition Al62(cubic, CsCl type, a = 2.9 A Cu25.6Fe12.4. In this case, the amount of b-phase was estimated to be lower than 5 wt%. The structural perfection of the quasicrystals can be estimated thanks to XRD peak broadening. Hence sample Al63Cu25Fe12 exhibits broad peaks and TEM has shown that this a modulated quasicrystal; quenching allows to obtain narrower peaks in the XRD diagram but the modulation is preserved. Other compositions exhibit narrow peaks by XRD and the structural quality of the quasicrystalline phase is good as observed by TEM
Fig. 1. SEM image from grains (white) belonging to the 25–53 lm particle size granulometry after crushing in an agate mortar.
Table 1 Compositions, samples preparation details and oxidation conditions Nominal composition (at.%)
Heat treatment
Phases identified before oxidation
Particle size (lm)
Oxidation conditions
Al63Cu25Fe12
800 C water quenched
Modulated icosahedral quasicrystal
<25
730 C water quenched
Modulated icosahedral quasicrystal
25–53 <25 25–53
500 C 500 C 500 C 500 C 500 C 500 C
Al62.5Cu25.3Fe12.2
730 C air quenched No heat treatment
Icosahedral quasicrystal Icosahedral quasicrystal
<25 <25
500 C in air 500 C in air
Al62Cu25.6Fe12.4
730 C air quenched
Icosahedral quasicrystal and b-phase.
<25
500 C in air
at at in in in in
1 · 104 mbar 1 · 106 mbar air air air air
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imaging and evidenced by the sharpness of the diffraction spots. Most samples were oxidized at 500 C in air using a conventional furnace. Powders were placed in an alumina container. In order to check the reproductibility of the experiment, it was done twice at two months interval for the sample having composition Al63Cu25Fe12. Moreover, some experiments were carried out under primary vacuum (1 · 104 mbar) and second– ary vacuum (1 · 106 mbar) in order to analyze the influence of the annealing conditions with and without an oxidatizing atmosphere. All samples were characterized at room temperature after cooling down from 500 C.
peak asymmetry and the amount of b-phase were then deciphered from the X-ray pattern. In order to determine the six-dimensional lattice parameter a6d we have used the indexing scheme proposed by Cahn et al. [8]. The AlCuFe icosahedral phase can be ascribed a Face Centered Cubic lattice in a sixdimensional real space with lattice parameter a6d. A simple relationship exists between inter-planar distances, lattice parameter a6d and diffusion vector qk: pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi sin h 2p N þ sM ˚ 1 Þ; pffiffiffiffiffiffiffiffiffiffiffiffiffi ðA ¼ jqk j ¼ 4p k a6d 2 þ 2s
A Siemens D500 diffractometer (Co anode, kKa1 = ˚ ) was used to characterize the powders. It is 1.78897 A fitted with a primary bent monochromator and a linear detector. The XRD patterns were recorded for 2h ranging from 25 to 120. The step size was equal to 0.032 2h, which corresponds to a channel of the detector, and the counting time was fixed to 1 h. In order to determine accurately the six-dimensional lattice parameter of ˚ ) was the icosahedral phase silicon powder (a = 5.4309 A used as an internal standard to obtain the experimental errors. Fig. 2 shows a powder X-ray diagram for the composition Al62.5Cu25.3Fe12.2 before oxidation. This diffraction pattern is representative of an icosahedral material of high lattice perfection. Diffraction peaks due to the internal silicon standard are also indicated. The X-ray diffraction patterns were fitted with an asymmetric Pearson VII profile shape function using the computer program Winfit [7]. The six-dimensional lattice parameter, FWHM (full width at half maximum),
where N and p Mffiffiffi are the indices of the diffracting planes and s ¼ ð1 þ 5Þ=2 is the golden number. The lattice parameterpffiffiffiffiffiffiffiffiffi is deduced from XRD peak ffi N þsM positions by plotting 2ppffiffiffiffiffiffiffi versus q k. One therefore 2þ2s obtains a straight line whose slope equals 1/a6d. The a6d parameter is a measure of the six-dimensional unit cell edge length. However, the three-dimensional quasiperiodic structure being obtained by a cut of 6D space, changes of a6d have the same meaning than in conventional crystallography. Thus, we can obtain information about the unit cell dimension and the atomic density, therefore about the formation of vacant sites, about structural changes due to enrichment or loss of an element (chemical change) and oxygen diffusion in the lattice. The relative changes in the FWHM were measured for peaks (70/113) and (72/116) as a function of oxidation time. As it is well known, peak broadening is linked with the occurrence of micro-strains and with crystallite sizes. However, for quasicrystals a relationship exists in the presence of a phason strain field between peak broadening and diffusion vector in perpendicular space q? that allows to quantify the strength of the phason field [9–12]. Fig. 3 shows the two peaks considered to obtain reliable data along the lines suggested.
Fig. 2. X-ray pattern for Al62.5Cu25.3Fe12.2 powder. Silicon peaks are pointed out (vertical bars) and indices for the main peaks of the icosahedral phase are given.
Fig. 3. Peaks (70/113) and (72/116) for a non-oxidized sample. The FWHM is close to the instrumental broadening.
2.2. X-ray diffraction characterization
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– and its lattice parameter for composition Al50Cu34Fe16 ˚ . This phase often coexists with the (at.%) is ab = 2.942 A i-AlCuFe phase because it forms during cooling from the liquid and participates to the peritectic reaction which causes the formation of the icosahedral phase. A ratio named F(b) was used as a marker to follow the growth of the b-phase. To define F(b), we used the intensities of the b-phase (1 1 0) peak and of the icosahedral (18/29) peak: F ðbÞ ¼ 100
Fig. 4. Icosahedral phase (72/116) peak used to quantify the peak asymmetry and (5 1 1) silicon peak.
This ratio does not give directly the b-phase volume fraction. However it can be calculated using the following formula: n¼
An asymmetry parameter a was defined as the difference between the half width at half maximum left (HWHMleft) and the half width at half maximum right (HWHMright): a = HWHMleft HWHMright (2h). This parameter has been calculated for peak (72/116). It is mainly due to a lattice parameter gradient, which can be inter-particles or intra-particles. If this gradient is only intra-particles, it can be due to oxygen or aluminum diffusion toward the grain surface. It must be emphasized that the penetration depth of X-rays at Co Ka wavelength in this AlCuFe alloy is around 20 lm. Fig. 4 shows the asymmetry of peak (72/116) of the icosahedral phase in an oxidized sample. One peak of the silicon standard is also shown. In one sample the b-phase (see Fig. 5 for indexation) is present previous to oxidation. It is a solid solution of composition Al50x(Cu,Fe)50+x which occupies a very broad field in the ternary system AlCuFe since nearly all iron can be substituted by copper. The b-phase structure belongs to the space group Pm 3m – CsCl structure
Ið1 1 0Þ . Ið18=29Þ
K Ið1 1 0Þ ; K Ið1 1 0Þ þ Ið18=29Þ
where n is the b-phase volume fraction and K is a constant determined by calibration on a sample containing 90%w of icosahedral phase and 10%w of b-phase powder (K 0.06).
3. Results The aim of this work was to study powder oxidation. However at 500 C, other mechanisms that would influence the evolution of various parameters can take place, namely phase transformation if equilibrium was not reached during cooling, crystallite growth and coalescence, etc. In order to study the effects of oxidation, we needed to make sure which changes are only due to oxidation. With this goal in mind, several samples were annealed at 500 C in a furnace under high vacuum (1 · 106 mbar). The evolution of lattice parameters in this condition is highly slowed down since the partial oxygen pressure is low. For all batches annealed in such conditions (composition Al63Cu25Fe12 at.%, annealed at 800 C, water quenched, particle size lower than 25 lm), a 178 h annealing causes a change in parameters, for instance a6d, similar to 2 h oxidation in air. After 178 h ˚ whereas annealing under vacuum, a6d = 6.3159 A ˚ a6d = 6.3156 A after 2 h in air. Thus it appears obvious that the parameter changes are mainly due to grain oxidation. 3.1. Evolution of the six-dimensional lattice parameter a6d
Fig. 5. Peaks positions and Miller indices for the main reflections of the b-phase: (1 0 0), (1 1 0), (2 1 1).
3.1.1. Composition effect Numerous papers dealing with quasicrystals give values for the six-dimensional lattice parameter but due to the various processing routes and characterization methods used, they can hardly be compared. However Quiquandon et al. [13] carried out an extensive study on
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the influence of the composition on the lattice parameter. They showed that the six-dimensional lattice parameter is proportional to the aluminum atomic content at constant electron/atom ratio (e/a). Using their experimental results we can infer a relationship between a6d, the Al at.% and the (e/a) ratio: a6d ¼ 0.3698 ðAl at:%Þ 0.1255 ðe=aÞ þ 6.32156. ð1Þ In our study the three selected compositions were chosen in order to keep the same ratio between copper and iron contents. Its is interesting to see that the lattice parameter is inversely proportional to the aluminum content. This behavior could be expected from Eq. (1) as can be seen in Fig. 6, however the slope is higher than expected and the difference increases with the aluminum content. This discrepancy could be due to the difference between the nominal composition of our samples and the real composition as well as to the presence of impurities as oxygen or carbon introduced during the sintering process. During oxidation the trend is the same whatever the composition, the initial heat treatment and the particle size. In all cases, the lattice parameter increases with oxidation time. This trend is shown in Fig. 7 for the three compositions, for particle sizes lower than 25 lm and for initially quenched samples. Using a linear time scale, the increasing lattice parameter in six-dimensional space seems to reach saturation beyond a few hundred hours. However, a6d evolution is controlled by atomic diffusion, which is governed by nonlinear equations. In this way, using a logarithmic time scale allows us to correct this effect. It appears that all series have a similar evolution: a rather linear increase of a6d with time in the beginning whereas it is only for long oxidation times
Fig. 7. a6d evolution with oxidation time at 500 C. The particle size is lower than 25 lm for the three series. Linear (a) and logarithmic time scale (b). Lines are only given to guide the eye. Error bars are typically of the same size than symbols.
that saturation comes in (higher than 1000 h of oxidation time). A slight decrease seems also to be observed for the largest oxidation times. 3.2. Particle size effect In order to understand the effect of particle sizes, the time dependence of the events is summarized in Fig. 8 in terms of lattice parameter for two different particle sizes. To check the reproducibility of the experiment, it was carried out twice for the Al63Cu25Fe12 batch. The lattice parameter increases more quickly for the lower particle sizes whereas the final lattice parameter is larger, which confirms that the process is linked with the specific area, because the effects of oxidation are more pronounced near the surface and are all the more amplified when the ratio area to volume is the largest. 3.3. Peaks width and asymmetry evolution
Fig. 6. Variation of the 6D lattice parameter a6d with the alloys aluminum content, before oxidation and comparison with Quiquandon [12] results. The continuous lines gives a6d using the equation: a6d = 0.3698 · (Al at.%) 0.1255 · (e/a) + 6.32156, derived from Quiquandon results.
The peaks full width at half maximum (FWHM) and asymmetry behave in a similar way as can be seen in Fig. 9. They both increase during the first step of oxidation, then reach a maximum and afterwards decrease to val-
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Fig. 8. Grain size effect on a6d evolution during oxidation at 500 C for powder Al63Cu25Fe12 – the second series of measurements was made several months after the first one.
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Fig. 10. Increase of the b-phase amount with oxidation time at 500 C for the composition Al63Cu25Fe12, particle size bellow 25 lm. Evaluation of the annealing temperature effect and of the reproducibility. The evolution of the (1 1 0) b-phase peak can be seen at the bottom of the graphic.
tion, whatever the composition may be, this phase systematically appears in amounts that increase with oxidation time as can been seen in Fig. 10 for several series. The richer in aluminum the alloy is, the later the b-phase appears and the lower the b-phase amount at saturation is. However, the various compositions follow all the same trend that is characterized by an increase at the beginning of oxidation and a saturation after an inflection point for longer oxidation times. The influence of the particle size on the b-phase amount can be assessed in Fig. 11. This variation is amplified when the particle size is the lowest, as it was already pointed out for the a6d parameter of the icosahedral phase. It must be emphasized that the lattice parameter of the b-phase does not change significantly with oxidation ˚ . FWHM and time and lies between 2.94 and 2.945 A asymmetry parameters exhibit similar behavior. From this result, it is straightforward to deduce that the bphase does not take part in the oxidation process, but is just an oxidation product. 3.5. Composition of the oxide layer
Fig. 9. FWHM (a) and asymmetry parameter a (b) evolution with oxidation time at 500 C for two particle sizes.
ues close to their initial values. Changes are more drastic for the smallest particle sizes. 3.4. Analysis of the b-phase Before oxidation, the b-phase is present in only one alloy, the poorest in aluminum. However during oxida-
The powders show a steel gray color that becomes darker and bluish with oxidation time. It is reasonable to assume that it is essentially aluminum, which oxidizes and iron in a lesser extent. Even for long oxidation times – three months at 500 C – no other crystalline phase than the b-phase and the icosahedral phase was detected. Particularly, no crystallized oxide appeared on the surface of the grains. From X-ray diffraction results it is not possible to conclude that the oxide layer is completely amorphous. Indeed the phase fraction could be too small to give rise to measurable peaks by XRD and diffraction peaks could be broadened by domain size effects. However a TEM study, whose results will be presented elsewhere, carried out on bulk samples, under the same oxidation conditions, showed that a
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Fig. 11. F(b) parameter variation with oxidation time at 500 C (a) for the two particle sizes for composition Al63Cu25Fe12 and (b) for various compositions (particle size bellow 25 lm). A logarithmic time scale is used.
homogenous and amorphous oxide layer appears on the sample surface during oxidation. This result is not clearly understood now because usually in aluminum alloys, amorphous aluminum oxide, which appears in the first step of the oxidation, quickly crystallizes for temperatures above 400–450 C [14].
asymmetry of the peaks can be due to a lattice parameter gradient and their values give information about the gradient intensity. The a6d parameter increases when the aluminum content decreases in the alloy. The same is observed when the b-phase appears and its amount increases. So, the b-phase appears when the loss of aluminum in the icosahedral phase is too significant, which is consistent with the phase diagram at 500 C. This aluminum loss is not uniform and even for the composition the richest in aluminum, it is likely that some zones quickly reach the 1 at.% loss in aluminum necessary to trigger the phase transformation. It is likely that the surface oxide layer is mainly composed of aluminum oxide with a small amount of iron in solution and no copper. The three compositions were chosen so that they stand on a straight line in the ternary phase diagram. During the oxidation process, the composition of the icosahedral phase probably follows this line, which was expected if only aluminum contributes the oxide layer. The plot of a6d with the b-phase ratio confirms this assumption since it can be seen in Fig. 12 that all compositions follow the same trend. This trend is linked with the [Cu]/[Fe] ratio and with the aluminum reservoir in the sample, i.e. with its specific area. This implies that for a given composition, the stability of the icosahedral phase depends on the quantity of aluminum exposed to contact with oxygen, as characterized by the surface area to volume ratio (S/V) of the particle. This effect is possible if aluminum diffusion is a long distance mechanism, at least over distances larger than 50 lm. If the diffusion was only a surface phenomenon, localized to within a few microns, no difference would be observed between different particle sizes. It must also be underlined that oxidation kinetics are similar whatever the size of the aluminum reservoir is. This kinetics slows down and seems to follow a parabolic evolution. After 1000 h oxidation, the oxidation rate is very low on a linear time scale and the phenomena seem to reach a saturation point.
4. Discussion 4.1. The aluminum reservoir effect In order to understand the evolution of the various parameters, it is interesting to point out the correlation that exists between each other. Whereas the width and the asymmetry of the diffraction peaks show similar evolution, the same is true for the a6d lattice parameter and the amount of b-phase, which by the way seem closely related. Moreover the first two parameters reach their maximum value at the half duration of the two latter parameters evolution, which corresponds to an inflection point. Thus, the increase of the width and of the
Fig. 12. a6d parameter evolution with b-phase fraction (F(b) parameter). Particle size bellow 25 lm.
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Fig. 14. Avrami plot for the sample Al62.5Cu25.3Fe12.2 (at.%), particle size <25 lm, oxidized at 500 C in air. Fig. 13. Plot showing the Al depletion in function of the oxide layer thickness for the sample Al63Cu25Fe12 (at.%) with four different particle sizes: 5, 25, 53 and 100 lm. The model uses spherical powder particles, and a homogeneous volumic depletion. The model also takes into account the atomic density differences (in Al at. nm3) between alumina and the quasicrystalline matrix. The dotted line at 62.2 at.% Al corresponds to the limit of the stability range of the quasicrystalline phase.
The Al depletion has been computed in function of the oxide layer thickness for different particle sizes (Fig. 13). Since the real S/V ratio could not be measured for the different samples, the model assumes the powder particles have a spherical shape, which is a rough approximation and leads to underestimate the Al depletion. Nonetheless, the diagram shows that for particle sizes larger than 100 lm it is unlikely that the icosahedral to b-phase transformation would occur with the oxidation conditions studied here.
Three processes pictured in Fig. 15 can be expected for the formation and growth mechanisms of the oxide. Each grain is embedded in a native oxide layer which forms during crushing. At 500 C, this layer grows following a parabolic kinetics because it implies aluminum diffusion toward the oxide/atmosphere interface. The diffusion coefficient being larger at the grain boundaries than in the bulk, it is likely than the aluminum depletion is more important at the boundary between two icosahedral grains and the oxide. Nucleation of the b-phase could possibly occur at such sites. The platelets could then grow from this interface towards the core of the icosahedral phase crystallites. Another process could lead to the formation of plates at metal/oxide interface by transformation of the icosahedral phase grains being
4.2. Kinetics of the b-phase transformation: an Avrami-law analysis Usually, an isothermal transformation can be analysed using the Avrami equation [15]: n ¼ 1 expðktn Þ; where n is the transformed volume fraction, k is the crystallization rate and the value of n is governed by the transformation mechanism. Christian [15] has given the various transformation mechanisms and corresponding n exponents. Thus, n can be determined using the following equation: 1 log log ¼ n logðtÞ þ logðkÞ. 1n The curve, plotted in Fig. 14 is obtained for the Al63Cu25Fe12 composition. The slope is roughly equal to 1/2, which indicates that the transformation is diffusion controlled, in accordance with the growth of broad platelets at the oxide/metal interface or/and at the grain boundaries where the aluminum depletion is the most significant.
Fig. 15. Schematic model for the three possible oxidation processes described in this paper. Powder grains are shown before, during the first oxidation steps and after a long oxidation time. (a) All crystallites of the grain transform into b-phase. (b) Only the crystallites near the surface of the grain undergo a phase transformation. (c) The transformation occurs at the grain boundaries and the b-phase crystallites look like platelets.
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in direct contact with the oxide layer. The last process, which is not incompatible with the two others, consists in the transformation of a whole grain of icosahedral in b-phase by growth at the grain boundaries or at the metal/oxide interface. The first process is the most plausible, but it is possible that the two nucleation processes coexist and that some grains fully transform in b-phase. 5. Conclusions We have shown that the oxidation behavior of i-AlCuFe powders at 500 C is mainly governed by diffusion of aluminum atoms towards the surface of the powder particles. Aluminum oxidizes and forms an amorphous oxide layer on the surface of the grains whatever the oxidation time. When the aluminum loss is too large, the icosahedral phase transforms into b-phase, in accordance with the phase diagram at 500 C. Another composition of the quasicrystal or another oxidation temperature would lead to a different phase transformation in accordance with the same phase diagram. Aluminum diffusion is a long distance process, which means that for a particle size larger than a critical particle size, appearance of the b-phase, due to a too important depletion in aluminum, can be avoided. This phenomenon was indeed observed in bulk samples, which have a lower specific area and contain a more important aluminum reservoir.
References [1] S.S. Kang, J.M. Dubois, J. Mater. Res. 10 (5) (1995) 1071. [2] D.J. Sordelet, M.J. Kramer, I.E. Anderson, M.F. Besser, in: C. Janot, R. Mosseri (Eds.), Proceedings of the 5th International Conference on Quasicrystals, World Scientific, Singapore, 1995, p. 778. [3] D.J. Sordelet, L.A. Guderman, M.F. Besser, A.B. Akinc, in: New Horizons in Quasicrystals: Research and Applications Conference Proceedings, Singapore, 1997, p. 296. [4] P. Weisbecker, G. Bonhomme, A. Cael, L. Zhang, J.M. Dubois, Mater. Res. Soc. Symp. Proc. K9 (2) (2001) 1. [5] B.I. Wehner, U. Ko¨ster, A. Ru¨diger, C. Pieper, D.J. Sordelet, Mater. Sci. Eng. A 294–296 (2000) 830. [6] M. Yamasaki, A.P. Tsai, J. Alloys Compd. 342 (2002) 473. [7] S. Kru¨mm, in: XIIIth, Conference on Clay Mineralogy and Petrology, Praha, 1994. [8] J.W. Cahn, D. Shechtman, D. Gratias, J. Mater. Res. 1 (1986) 13. [9] M. Jono, Y. Matsuo, Y. Ishii, Mater. Sci. Eng. A 294–296 (2000) 680. [10] M. de Boissieu, H. Takakura, M. Bletry, J.Q. Guo, A.P. Tsai, J. Alloys Compd. 342 (1–2) (2002) 265. [11] A. Le´toublon, I.R. Fisher, T.J. Sato, M. de Boissieu, M. Boudard, S. Agliozzo, L. Mancini, J. Gastaldi, P.C. Canfield, A.I. Goldman, A.P. Tsai, Mater. Sci. Eng. A 294–296 (2000) 127. [12] A.P. Tsai, A. Inoue, T. Masumoto, Prog. Cryst. Growth Char. 34 (1997) 221. [13] M. Quiquandon, A. Quivy, J. Devaud, F. Faudot, S. Lefebvre, M. Bessie`re, Y. Calvayrac, J. Phys.: Condens. Matter 8 (1996) 2487. [14] K. Wefers, C. Misra, Oxides and hydroxides of aluminum, Alcoa Technical Paper no. 19, Alcoa Laboratories, 1987. [15] J.W. Christian, The Theory of Transformations in Metals and Alloys, Pergamon, London, 1965.