CRYSTALLINE AND QUASI-CRYSTALLINE STRUCTURES IN AN Al-Li-Cu-Mg ALLOY

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CRYSTALLINE AND QUASI-CRYSTALLINE STRUCTURES IN AN Al-Li-Cu-Mg ALLOY

A. Loiseau, G. Lapasset

To cite this version:

A. Loiseau, G. Lapasset. CRYSTALLINE AND QUASI-CRYSTALLINE STRUCTURES IN AN Al-Li-Cu-Mg ALLOY. Journal de Physique Colloques, 1986, 47 (C3), pp.C3-331-C3-340.

�10.1051/jphyscol:1986334�. �jpa-00225746�

JOURNAL DE PHYSIQUE

Colloque C3, supplgment a u n o 7, T o m e 47, juillet 1986

CRYSTALLINE AND QUASI-CRYSTALLINE STRUCTURES I N A N Al-Li-Cu-Mg ALLOY

A. LOISEAU a n d G. LAPASSET

ONERA, 29 Av. de la Division Leclerc, BP. 72, F-92322 Chatillon, France

Abstract - We have studied by means of transplission electron microscopy the particular properties of an icosahedral phase (I-phase) formed in equilibrium conditions in an A1-Li-Cu-Mg alloy. For instance this I-phase is observed as long faceted precipitates embedded in the f.c.c.

A1 matrix and exhibits definite orientation relationships which were un- expected considering the natural orientation between a cube and an ico- sahedron. The I-phase is observed together with a crystalline phase (C-phase) which is strongly related to it, since it gives pseudo five- fold symmetry diffraction patterns fitting almost perfectly the five- fold symmetry pattern of the I-phase. First results on this unknown phase are reported.

1 - INTRODUCTION-

The quasi-crystalline phase (I-phase) which exhibits a long range orien- tational order with icosahedral point group symmetry m35, has been discovered in the Al-14at%Pin alloy /1/ and has been afterwards observed in various

A1-10-20at%T alloys /2,3/ where T is a transition metal (Re, Cr, Mo, V, Ru, Fe, Co, Ni) and also in (Ti,V) Ni /4/, U-Pd-Si

Al-Cu-Mg

In all these cases, an important point is that the I-phase can only be obtained by rapid so- lidification techniques (melt-spinning for instance). The I-phase has generally a dendritic morphology and can coexist with the amorphous phase or with a crys- talline phase (f.c.c. Al). Urban et a1 171 have shown in A1-)In and A1-V that the I-phase fully transforms into the amorphous phase under electron irradiation at low temperature; when heating the sample, the amorphous phase transforms into tho I-phase which transforms then itself into a crystalline phase at higher temperature.

However the I-phase can be formed in entirely different conditions.

Sainfort et a1

have indeed observed it in A1-Li-Cu and Al-Li-Cu-Mg alloys which are produced by classical casting methods. In this case, the I-phase is formed in as-equilibrium conditions by solid state precipitation in the homo- genized solid solution during thermal treatments in a wide range of temperature.

In these alloys, due to its formation mechanism, the I-phase exhibits quite original properties such as its morphology or orientation relationships with the f.c.c. matrix. Furthermore, besides the I-phase, these alloys can contain a crystalline phase which is strongly related to it. In the A1-Li-Cu al- loys 191, this phase is the R phase (point group symmetry Im3) which is isomor- phous to the a -AlMnSi phase: in this latter phase, the atomic structural unit is built with Mn icosahedrons, each of them containing an A1 icosahedron. On this basis, Audier et a1 /10,11/, Henley et a1 112,131 have proposed structural models for the I-phase. In these models, the a phase can be seen as an "appro- ximant structure" of the I-phase in the sense defined by Henley et a1 1121 and Gratias et a1 1141. The R phase is unfortunately not observed in the A1-Li-Cu-Mg alloy we have studied but it is replaced by another unknown phase which is denoted C-phase in the following.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1986334

In this paper, we report on a study by means of electron microscopy of the structure of the I-phase and of the related crystalline phase in a commer- cial Al-Li-Cu-Mg alloy. Results obtained in the ternary Al-Li-Cu system by Sainfort et a1 191 are detailed elsewhere in this issue.,

2 - EXPERIMENTAL PROCEDURE -

The alloy is a 8090 Aluminium alloy. Weight percentages of addition ele- ments are respectively Li 2.5%, Cu 1.3%, Mg 1%. This alloy was elaborated by CEGEDUR PECHZNEY using ingot-casting methods. After the solution heat treatment in the monophase field, the samples are water-quenched and then annealed at 400°C for different times (lh., 4h., 24h., 32 days). As a consequence of these thermal treatments, special care was brought to remove Li- and Mg-depleted sur- face layers. Thin foils for TEM investigations were double-jet electrothinned in a 113 HNO - 213 methanol electrolyte at -40°C.

3 3 - ^RESULTS -

Four major phases are observed for all annealing times:

- f.c.c. A1 solid solution

- ^6'(L12 A13Li phase): the demixtion of the 400°C

solid solution during the cooling and at room temperature leads to a fine precipita- tion which gives rise to weak surstructure spots

(100) and (110)

in the A1 reciprocal lattice.

- an I-phase, which nucleates at the grains and subgrains boundaries as well as in the Al matrix. This phase appears from the beginning of the annealing and then coarsens. After the 24 h. annealing time, the size of the precipitates does not evolve significantly and is roughly a few micro- meters.

- an unknown crystalline phase (C-phase). This phase preci- pitates at the grains and subgrains boundaries and undergoes a heavy coar- sening. Scanning electron microscopy examinations (fig. 1) show ,for a 24 h.

annealing, large particles (up to 20 micrometers) of the C-phase as imaged by backscattered (fig. la) or secondary electrons (fig. lb). Precipitates of the I-phase are responsible for small bright dots in fig. la which disappear in fig. lb.

It must be emphasized that these two last phases are not observed in the as-quenched state i.e. before annealing at 400°C. The most striking point is their coexistence at 400°C: nevertheless, it is not yet clear whether these phases are equilibrium phases or whether one phase transforms into the other during the annealing.

figure

SEM image of the A1-2.5Li-1.3Cu-1Mg alloy after a 24l-i. annealing at 400°C.

a) image formed with backscattered electrons.

b) image formed with secondary electrons.

3 - 1 Results on the I-phase.

This phase displays the typical and now well-known electron diffraction

patterns of an icosahedral phase. Figure 2 shows the diffraction patterns along

the two-fold (A2). the three-fold (A3) and the five-fold (A5) axis. which are

quite identical to those observed in A1-Mn alloys 11,151 provided an homothetic factor is introduced. Angles between the A2, A3 and A5 axis are those expected for the m33 group symmetry. The figure 3 is a high resolution image obtained along the 5-fold axis, which reveals typical aperiodic arrangements of penta- gons of white dots. As pointed out by Sainfort et a1 for the A1-Li-Cu alloys 191, misalignements of diffraction spots along the two-fold directions are so- metimes observed in our alloy. For instance, misalignements can be seen in the three-fold symmetry pattern of fig 2b while they are not present in the five- fold symmetry pattern of fig 2c which is obtained on a different precipitate.

This misalignement effect has been reported in melt-quenched A1-Mn alloys by

figure 2 : typical diffraction patterns of the I-phase : a) 2-fold symmetry pattern b) 3-fold symmetry pattern c) 5-fold symmetry pattern.

figure 3

high resolution image of the I-phase along a 5-fold axis (R. Portier CNRS-Vi try

some authors /7,16,17/, but has never been observed in Al-Mn-Si alloys 1181.

Dixit et a1 /17/ have shown that the displacements disappear after annealing,

so they attribute them to structural defects. This is probably also the case in

A1-Li-Cu-(Mg) alloys. Further work is currently in progress in order to deter-

mine the origin and the nature of these defects in our alloy.

The intragranular I-phase precipitates exhibit three definite orienta- tion relationships with the A1 matrix. As it can be seen further, these orien- tation relationships are very different from the simple one Sainfort et a1 /9/

have found in the case of the Al-Li-Cu system. In this orientation relation- ship, a set of orthogonal A2 axis is associated with the set of<100>A1 axis and the<lll>Al directions coincide with four A3 I-phase axis: this charac- terizes the natural orientation between a cube and an icosahedron (see fig 3 in 191 1-

In our 8090 alloy, the three orientation relationships hold a common feature: an A2 axis is parallel to an<llO >A1 axis. Moreover, for any precipi- tate, the diffraction pattern along this basic A2 axis allows a complete crys- tallographic description of the precipitate and of the matrix: this is illus- trated in fig 4 successively for each of the three orientation relationships.

The drawn rhombi are useful to describe the two A2 directions (the rhombus dia- gonals) and the two A5 directions (the rhombus edges) which are perpendicular to the basic A2 axis. Resulting stereographic projections are given in fig 5.

In the major orientation relationship (a very rough estimation of its frequency is 80%), the two A2 directions (normal to-the basic A2 axis or [Oll]

A1 axis) are respectively parallel to [Ill] and [211] A1 axis (fig 4a). Dif- fraction patterns along these two particular directions are shown successively in fig 6a and 6b. Furthermore, an A5 axis (perpendicular to the basic A2 axis) lies near an [loo] axis (2" off).

The second orientation relationship is illustrated in fig 4b: the two perpendicular A2 and the two A5 axis (all normal to the basic axis) arz pa- rallel or nearly parallel respectively to the [loo], [Oll], [Ill] and [Ill] A1 directions.

In the case of the third orientation relationship, one of two these A5 axis (fig 4c) coincides exactly with the [TIT] A1 direction (see the (111) dif- fraction pattern-in fig 7). An A3 axis (perpendicular to the basic A2 axis) is close to the [Oll] A1 direction.

All of these relationships bring an A5 axis near a<1127A1 direction.

Fig 8 shows the diffraction pattern along this direction for the first orienta- tion relationship.

figure 4

a) b) c) the 3 different orientations between 2-fold symmetry pattern of the I-phase and the (011) A1 diffraction pattern allowing complete determination of the orientation relationships. Strongest I-phase spots build a rhombus enhanced with white lines: the rhombus diagonals are along 2-fold axis and its edges are along 5-fold axis.

d) is the (011) zone of the matrix alone: the weak surstructure

spots (100) are due to the fine precipitation of the AI3Li (L12)

phase.

figure 5

the three orientation relationships - labelled I,

- observed between the I-phase and the matrix as defined by the relative rotation between the stereographic projection of t.he I-phase along the basic 2-fold axis and the projection of the matrix along the

[Oil] axis parallel to this A2 axis. Only the f.c.c. poles lying in the (011) plane are sketched.

figure 6

diffragtion patterns typical of the orientation relationship I:

a) (111) Al zone axis, parallel to a 2-fold I-phase axis.

figure 7

diffraction pattern typical of the third orientation relationship:

(Tli) A1 zone axis is exactly parallel to a 5-fold I-phase axis.

figure 8

a common feature typical of the three orientation relationships:

an (112) A1 zone axis is nearly parallel to a 5-fold I-phase axis.

The I-phase precipitates have a "pencilw-like morphology, with the basic A2 direction as the pencil axis. In fig 9a ((111) A1 zone axis)2 several preci- pirates can be seen their growth axis being parallel to the [Oll], [10T] and

[I101 Al directions. These precipitates are faceted as it is clearly visible in fig 9b where the "pencil" axis is parallel to the electron beam. In the case of the first orientation relationship, the most developped facet is normal to [loo] A1 i.e. to an A5 axis.

figure 9

bright field images of the I-phase pre- cipitates in different orientations revealing their faceted pencil shape.

a) the plane of the figure,nomal to (111) Al, contains three < 110> direc- tions at 60° each other,corresponding to three orientations of growth axis of the precipitates.

b) the growth axis is normal to the fi- gure: the faceting is clearly visible.

C)

and d) the plane of the figure con- tains the growth axis: a well-develop- ped facet is edge on in d) and is lying in the figure plane in c).

3 - 2 Results on the related crystalline phase (C-phase)

The structure of the C-phase is not yet resolved. We emphasize here its main structural characteristics and its relations to the I-phase.

Microprobe analysis of very coarse C-phase particles gives Cu and Mg enrichment compared to the matrix composition. Assuming that this phase is a Li-bearing compound, its composition is (at%): A1 58.3% (r 1.6%), Cu 5.0%

0.4%), Mg 4.2%

0.2%) and Li 32.5% (T 1.8%). The Cu content is remarkably lower in the C-phase than in the R-A1 CuLi3phase. Moreover, in the ternary sys- tem, the Cu content is higher in the i-phase than in the T2-Al6CuLi3 phase (i.e. the I-phase); it seems to be exactly the contrary in our alloy where, as estimated by X-ray energy dispersive measurements in the electron microscope, the Cu content is higher in the I-phase than in the C-phase.

The structure-of the ~ - ~ h a s e is strictly periodic wf th tetragonal point

group symmetry. One of the most interesting points is that certain electron

diffraction patterns obtained in this phase exhibit very strong similarities

with the five-fold symmetry pattern of the I-phase, as shown in fig 10. Fig lob

corresponds to the C-phase: a periodic array of spots is clearly seen. By

considering now the distribution of the most intense spots, which is due to the

symmetry of the motif, one can distinguish concentric slightly distorted rings

of ten spots. The angles between the spots are not exactly

36O but vary

between 34 and 38" i.e. are at most 2 ' off the exact value. It is interesting

to compare these rings with those observed for the I-phase in the fivefold sym-

metry pattern (fig 10a). Both patterns have been superimposed in fig 10c where

the white dots represent the I-phase and the black dots the C-phase. In the

figure 10: a) 5-fold symmetry pattern of the I-phase.

b) pseudo 5-fold symmetry pattern of the C-phase: the most intense spots are distributed on successive slightly distorted rings of ten spots.

c) superposition of both patterns

I-phase spots; and e: C-phase spots). Drawn lines and rings correspond to the I-phase

I-phase, the spots issued from the successive rings are perfectly aligned while in the C-phase this alignement occurs only in one direction (direction

which can be indexed as a [hhO] C direction). In the other directions, the position of the spots deviates from the perfect one and oscillates around it, fitting it better and better as the distance increases. Considering the direction of per- fect alignement, the successive positions of the intense spots are in ratio 2/1,

8/5 which are the first approximants of the golden mean. This remarkable succession of the intense spots can also be seen in fig 11 which is the four-fold symetry pattern ((00E) C zone axis) containing the [hhO]C di- rection and its equivalent.

The striking properties of the diffraction pattern shown in fig lob which has a pseudo five-fold symmetry is an evidence of a local icosahedral symmetry of the atomic basic unit. It is however very difficult to recognize other patterns which could exhibit pseudo two-fold or three-fold symmetry simi- lar to those of the I-phase. Fig 12 shows on the standard stereographic

triangle the various observed diffraction patterns (DP a to k). Referring to

the stereographic triangle in the inset, one can index all diffraction pat-

terns. UP a to g contain the [h60] C row. The DP c has the pseudo five-fold

symmetry. The DP a is the four-fold symmetry pattern ((0OR)C zone axis).

figure 1 1

(0011) four-fold symmetry pattern of the C-phase. Along the [hhOl row the succes- sive positions of intense spots are in ratio 211, 312, 513, and 815.

The DP h and g correspond respectively to the (OhO) and to the (hhO) C zones: they exhibit the second important characteristic of the C-phase. Along the [OOR] C row, the distance between two consecutive reflections is very short, indicating the existence of a very long period along the c axis. Fig 13 is a bright field image of a precipitate obtained from one of these orienta- tions: the selecting aperture used includes some satellites reflections in ad- dition to the transmitted beam, so that an array of thin interference fringes appears. The average distance between the fringes is about 7 nm which corres- ponds to the shortest distance between two spots along the [OOR] row. Now, considering only the most intense spots in the (OhO) DP (DP h), they form a four-fold symmetry array which fits quite perfectly with that of the (0011) zone (DP a).

These results allow us to conclude that the C-phase is based on a rough- ly cubic structure with a parameter a 1 2 nm (five times the A1 matrix para- meter), with a modulation of planar defects along a cube axis which has a very long period about 7 nm. These planar defects can be stacking faults as anti- phase boundaries or twins. Obtention of high resolution image of this structure is therefore necessary to elucidate the nature of these defects. This work is currently in progress. The nature of the modulation has to be compared with that observed in the T-phase discovered by Bendersky 1191. Some patterns of both C-phase and T-phase are indeed streaking similar. However, the T-phase has an exact decagonal symmetry and is not a crystalline phase as the C-phase.

It should be mentioned that Ball et a1 1201 have briefly described four types of precipitates including the I-phase (type 4 precipitates) which have been observed in an A1-2.5 Li-1.2 Cu-0.7 Mg (wt%) alloy. They reported for the first and the second types of precipitates diffraction patterns which all have been found to belong to the C-phase. These two types of precipitates may in fact correspond to a single phase, the C-phase.

Structural connections between the C-phase and the I-phase are not yet clear since the structure of the C-phase is actually not solved. This structure contains a motif with a local pseudo five-fold symmetry, as well as in the

a-AlMnSi phase. However, the C-phase cannot be seen as an approximant struc-

ture of the I-phase in the same sense as the a-phase. This difficulty comes

from the fact that the C-phase has a tetragonal symmetry instead of a cubic

one: the set of the three orthogonal<100> axis does not correspond to a set of

three orthogonal A2 axis in the I-phase, which is the basis of the transforma-

tion of the I-phase into an approximant structure like the a-phase. One may

correlate this difficulty with the unexpected nature of the orientation rela-

tionships between I-phase and f.c.c. matrix observed in our alloy. In addition,

we have noted that the long axis of the I-precipitates is along a two-fold axis

parallel to a [110] Al axis. C-precipitates can have also an elongated shape

with well developped facets and with a long axis parallel to a [hhO] C axis and

to an I-phase A2 axis.

figure 12: various diffraction patterns of the C-phase.

refer to the standard stereographic triangle put in the inset for indexing the pattern. Note in particular: a) (001) 4-fold symmetry pattern

C )

(111) zone g) (110) zone h) (010) zone i) (011) zone. c) is

the pseudo five-fold symmetry pattern.

figure 13: bright field image of the C-phase showing fine fringes due to a modulation of planar defects along the c axis.

ACKNOWLEDGEMENTS

The authors wish to thank CEGEDUR-PECHINEY Voreppe Research Center for providing the alloy and R. Portier from CNRS Vitry for the H.R.E.M. observa- tions. They are also grateful to D. Gratias (CNRS Vitry) and to F. Ducastelle (ONERA) for helpful discussions.

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