Metastable phases in mechanically alloyed AlMn powder mixtures

Materials Science and Engineering, A 131 ( 1991 ) 237-242

237

Metastable Phases in Mechanically Alloyed AI-Mn Powder Mixtures C. SURYANARAYANA* and R. SUNDARESANt Wright Research and Development Center, WRDC/MLLS, Wright-Patterson Air Force Base, OH 45433-6533 (U.S.A.) (Received April 27, 1990; in revised form June 19, 1990)

Abstract Elemental aluminum and manganese powder mixtures have been mechanically alloyed in a Spex mill. It has been shown that supersaturated solid solutions of aluminum containing as much as 18.5 at. % Mn can be produced in the as-milled condition. Room temperature aging produced a metastable f.c.c, phase with a =0.4472 nm. The equilibrium AloMn phase formed by decomposition of the supersaturated solid solution on elevated temperature exposure of the powder at temperatures higher than 623 K. The activation energy for decomposition has been calculated to be 2. 04 e V, the same as the activation energy for the diffusion of manganese atoms in an aluminum matrix.

1. Introduction Aluminum-manganese alloys are technologically important because of their high specific strength, modulus of elasticity and superior corrosion resistance in addition to their easy fabricability. In spite of this, the A1-Mn phase diagram has not been accurately determined till recently [1 ]. A critical evaluation of the phase diagram, based on both theoretical and experimental approaches, revealed the presence of several stable and metastable phases in the A1-Mn system [1]. Further, it has been shown that quasicrystalline icosahedral [2] and decagonal [3] phases with the forbidden fivefold and 10-fold rotational symmetries respectively can also be produced in aluminum-rich AI-Mn alloys by rapid quenching from the liquid state. In fact, in recent years AI-Mn has become the model system to verify the formation of quasi-crystalline *Permanent address: Department of Metallurgical Engineering~Banaras Hindu University, Varanasi 221005, India. tPresent address: Defence Metallurgical Research Laboratory, Hyderabad 500258, India. 0921-5093/9 !/$3.50

phases by a variety of non-equilibrium processing techniques [4-6]. The present investigation was undertaken with a view to produce metastable phases through mechanical alloying, and the present contribution reports on the formation of a supersaturated solid solution of manganese in aluminum and a new metastable f.c.c, phase in mechanically alloyed aluminum and manganese powder mixtures. Mechanical alloying is a powder-processing technique and involves the intimate mixing of powders, during which process novel materials with improved properties are produced by the competing actions of cold welding, fracture and rewelding of the powder particles. It has been suggested that mechanical alloying is applicable to the production of true alloys even at low temperatures owing to the high energies involved in ball milling [7], although examples of true alloying achieved are limited [8]. This process results in the formation of metastable supersaturated solid solutions [9], crystalline intermediate phases [10] and metallic glasses [11, 12]. Another useful and important attribute of mechanical alloying is the development of extremely fine grain sizes reaching down to nanometer levels [13] and a very uniform distribution of dispersoids which resist coarsening up to reasonably high temperatures [14]. 2. Experimental procedure Pure elemental aluminum and manganese powders were mixed to yield an average starting composition of 33.7 wt.% (20 at.%) Mn. This composition was chosen because most of the quasi-crystalline phases were produced homogeneously at this composition. Mechanical alloying was carried out at room temperature in a Spex 8000 Mixer Mill for times ranging from 12 to 36 h. The grinding medium was hardened 52 100 steel balls of diameter f~ in. About 1% stearic acid © Elsevier Sequoia/Printed in The Netherlands

238

was also added as a process control agent. For each run, 10 g of the powder and approximately 40 g of steel balls were loaded into the canister in an argon-filled glove-box. Forced air cooling during milling prevented excessive rise of the temperature of the powder. Further, to improve intimate mixing and to avoid formation of a thick crust at the bottom, the canister was rotated through 90 ° every 30 min and was inverted (through 180 °) every 2 h. The mechanically alloyed powder was removed from the canister and subjected to X-ray diffraction using monochromatic C u K a radiation at 40 kV and 100 mA settings in a Rigaku X-ray diffractometer. The phases present were identified by comparing the peak positions and intensities with those listed in the JCPDS files. The mechanically alloyed powder was also annealed in vacuum-sealed silica capsules for 1 h each at different temperatures to follow the changes in the nature, identity and proportion of the phases. Differential scanning calorimetry (DSC) of the powders was also carried out to follow the transformation temperatures at different heating rates and to calculate the activation energies for the decomposition process.

in any change in the situation and only these two phases continued to be present. Since the extent of alloying was very limited in the as-milled powder, it was decided to investigate the structural changes, if any, taking place in the powder mix as a function of temperature. With this purpose in view, the as-milled powder mixture was annealed for 1 h each at 623 and 723 K and quenched to room temperature. X-ray diffraction patterns of these annealed samples indicated that in addition to f.c.c, aluminum and cubic a-Mn phases, the equilibrium AI6Mn intermetallic compound with an orthorhombic structure and the lattice parameters a = 0.64978 nm, b--0.75518 nm and c = 0 . 8 8 7 0 3 nm had also formed (Fig. 2). It has also been noted that the amount of the AI6Mn phase increased with an • increase in the temperature of annealing. For example, the volume per cent of the AI6Mn phase was only about 6% at 623 K but increased t o about 21% on annealing at 723 K (Table 1 ). Aging of the as-milled powder mixture at room temperature for 30 days resulted in the observation of some extra peaks in the X-ray pattern (Fig. 3). These extra peaks could be identified as

SSSO

3. Results

Figure 1 shows the X-ray diffraction pattern of the AI-Mn powder mixture milled for 12 h. An analysis of the patternindicated that only f.c.c. aluminum and cubic a-Mn phases were present in this condition. Mechanical alloying for even longer periods of time (up to 36 h) did not result

2907

2408

1

~

20a2

le40 1|17

4820[

GO 2 THETA 80

I 120

100

Fig. 2. X-ray diffraction pattern of the mechanically alloyed powder annealed for 1 h at 723 K. The pattern indicates the presence of aluminum solid solution, a-Mn and the equilibrium orthorhombic AI6Mn phases.

3888

~ 1928

TABLE 1 Relative proportions of the phases present in mechanically alloyed AI-Mn powder mixtures

ge4

oJso

40

Condition 4~

*'6 2 THETA e',

,~2

,~o

Fig. 1. X-ray diffraction pattern of the AI-Mn powder mix mechanically alloyed for 12 h. Only aluminum solid solution and a-Mn phases are observed in this condition.

AS milled Annealed 623 K, 1 h Annealed 723 K, 1 h

Relative proportion (vol.%) AI.

Mn

AI~Mn

98.9 91.6 76.5

1.1 2.7 2.5

-5.7 21

239

arising from an f.c.c, phase with a = 0.4472 nm. Table 2 lists the indexing of these peaks and it can be seen that the matching between the observed and calculated interplanar spacings is very good. On further aging of the same powder mixture for an additional 20 days, the peaks corresponding to the new f.c.c, phase disappeared. Thus this f.c.c. phase can be considered metastable. It may also be pointed out that even though only f.c.c, aluminum and cubic a-Mn phases were observed after 50 days of aging, in comparison to the as-milled

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powder the relative intensity of the aluminum peaks decreased and that of a-Mn increased in the aged powder. DSC investigations have been carried out to follow the decomposition behavior of the metastable phases produced in the mechanically alloyed powder mix. An exothermic peak was observed at 705 K on heating the powder mix at a rate of 5 °C min-1. The activation energy for the transformation was determined by noting the changes in the transformation peak temperature Tp with different heating rates ft. These two parameters are related to each other through the activation energy for the process, E, by the equation log(Tp2/fl)ocE/Tp [15]. Values of Tp were determined at heating rates ranging from 5 to 40°C min -l. A plot of log(Tp2/fl) against 1/Tp gave a straight line (Fig. 4), from the slope of which an activation energy for the process has been calculated as 47 kcal mol- ~(2.04 eV ).

2864

4. Discussion 77

2 THETA

Fig. 3. X - r a y d i f f r a c t i o n p a t t e r n o f t h e a g e d f o r 30 d a y s at r o o m t e m p e r a t u r e . I n n u m solid s o l u t i o n a n d t h e a - M n p h a s e s , p h a s e also a p p e a r s at this stage. S e e T a b l e of t h e p h a s e s .

96

as-milled powder a d d i t i o n to a l u m i a m e t a s t a b l e f.c.c. 2 for the indexing

Mechanical alloying has been known to be a powerful technique for producing metastable phases in a variety of alloy systems [16]. In many cases the metastable phases produced are very similar to those produced by rapid solidification from the liquid state, although the composition

TABLE 2 12 h ( M A )

Indexing of the X-ray difl'raction patterns of aluminum and manganese powder mixtures mechanically alloyed for and aged for 30 and 50 days

Line no.

d,~b, ( n m )

MA ~

l 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17

I).2342 -0.2106 0.2025 0.1908 11.1822 0.1750 -0.1432 --0.1220 0.1169 -0.0928 0.0905

~Fig. 1. bFig. 3. CMetastable p h a s e .

Al,~, f.c.c, a = 0.4032 n m M A + 30 days aging b

M A + 50 days aging

11.2584 0.2328 0.2233 0.2095 0.2018 0.1895 . . 0.1582 0.1426 0.1351 0.1289 0.1216 0.1165 0.0951) 0.0923 --

. 0.2325 . 0.2105 0.2025 0.1904 . . . 0.1430 . . 0.1221 0.1170 . -0.0905

. .

dcaI

. .

ct-Mn, c u b i c a = 0.8913 n m hkl

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.

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.

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.

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111 .

.

.

.

0.2101 . . 0.1900 0.1819 0.1748

330 . 332 422 431

. -200 --

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F.c.c5 a = 0.4472 n m

. 220

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.

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. 331 420

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0.2582 . 0.2236 -. ---0.1581 . 0.1348 0.1291 . . 0.0913 . .

l l 1 200 ----220 311 222

422

240 5.0

A

4.8"

~ ' ~ 4.6 4.4 .a

4.2

~1

4.0 1.32

I I " I mI m

1.34

1.36 !.38

1.40

1.42

I / T p x l 0 3 ( K "l) Fig. 4. Kissinger plot to calculate the activation energy of the process.

ranges over which these phases are formed are different in the two methods, largely owing to the differences in the mechanisms of their formation [17]. Quasi-crystalline icosahedral and decagonal phases have been reported to form homogeneously in AI-Mn alloys containing about 20 at.% Mn by a variety of processing techniques, including rapid solidification, vapor quenching, electrodeposition, electron beam surface melting, radio frequency ion plating, rapid pressure application, gas evaporation, irradiation of the amorphous phase, interdiffusion of evaporated layers and precipitation from a supersaturated solid solution [4-6]. Since mechanical alloying is another proven non-equilibrium processing technique, it has been recently reported that quasi-crystalline phases were produced in A1-Mn [18] and M g - A I - Z n and Mg-AI-Cu alloys [19] through this technique also. In our investigations, however, the X-ray diffraction patterns recorded from the powder in the as-milled condition revealed that no quasi-crystalline phase had formed. Transmission electron microscopy and diffraction experiments also confirmed this observation. Even more surprising is the fact that even after 36 h of high energy ball milling, only f.c.c, aluminum and a-Mn phases were observed, suggesting apparently that at room temperature there is no alloying even to produce an equilibrium intermetallic phase. Elevated temperature exposure of the powder, associated with increased diffusivity of the species, led to the formation of the equilibrium AI6Mn phase.

4.1. Supersaturated sofid solution The above results can be rationalized by considering the phase proportions in the as-milled powder mix. The volume fraction of only 1.1% Mn suggests that a considerable amount of man-

ganese has gone into solid solution in aluminum. Applying the lever rule to this two-phase mixture, one can calculate the manganese content in the aluminum solid solution phase and it turns out to be 31.7 wt.% (or 18.5 at,%). From this observation it becomes clear that mechanical alloying has resulted in a substantial increase in the (metastable) solid solubility of manganese in aluminum. Although increased solubility has been observed by mechanical alloying in some alloy systems earlier, e.g. 5 at.% Zr in iron [9, 17], 7 at.% Ni in zirconium [20] and 6 at.% Mg in titanium [13], this is perhaps the first time that such a large extension is reported. This large amount of supersaturation, coupled with the frequently made observation that kinetic restrictions of the mechanical alloying process prevent the formation of equilibrium intermetallic phases [21], can account for the absence of the equilibrium A16Mn or other intermetallic phases in this system on room temperature mechanical alloying. It is natural to expect that such a large solid solubility should lead to a change in the lattice parameter of the aluminum solid solution. However, the atomic radii of aluminum (0.182 nm) and manganese (0.179 nm) are so close to each other that there is only a slight decrease in the lattice parameter of aluminum solid solution under equilibrium conditions [22]. Accordingly, the aluminum solid solution in the mechanically alloyed condition has a lattice parameter of 0.4032 nm, only slightly smaller than the value of 0.4049 nm reported for pure aluminum. Elevated tempeature exposure of the powder mix containing the aluminum solid solution and the a-Mn phases resulted in the decomposition of the supersaturated solid solution leading to the formation of AI6Mn. As expected, the amount of the AI6Mn phase increased with increasing temperature and the volume fraction was as high as 0.21 at 723 K. It is anticipated that higher temperature annealing will result in further progress of alloying with complete elimination of free manganese.

4.2. Metastable fc.c. phase It was mentioned that aging of the powder mix containing the aluminum solid solution and the cubic a-Mn phases for 30 days at room temperature resulted in the formation of a metastable f.c.c, phase. Since the aluminum supersaturated solid solution containing manganese is already in a highly metastable state, formation of metastable

241

transition phases during decomposition is not uncommon. Thus the metastable f.c.c, phase would have formed to lower the free energy of the system and can be considered a precursor to the formation of the equilibrium A16Mn phase, which was achieved only on exposure of the powder to elevated temperatures of the order of 623 K and higher. The presumption that the decomposition of the supersaturated solid solution led to the occurrence of the metastable f.c.c, phase is further strengthened by the observation that the intensity of the aluminum solid solution lines in the X-ray diffraction patterns decreased on further aging and approaching the equilibrium. However, the only point of concern here is that the "stability" of this metastable phase is really low, disappearing just after 20 days of aging at room temperature. Further, the mechanism for the formation of this metastable phase is not clear at present. It can, however, be said with confidence that this is not impurity stabilized (due to oxygen, carbon, etc.), because on further aging, this metastable phase has completely transformed to the equilibrium phases. The activation energy determined from the peak shift method of Kissinger for the decomposition of the supersaturated solid solution and/ or the metastable f.c.c, intermediate phase is 47 kcal mol J (2.04 eV). This value is significantly lower than that for the crystallization of metallic glasses produced by rapid solidification from the melt ]23], suggesting again that the thermal stability of the metastable phase produced in the mechanical alloying process is not very high. In this context it may also be noted that the activation energy for the diffusion of manganese in aluminum has been reported to be 1.96 eV [24], 2.19 eV [25] and 2.25 eV [26], very close to the value calculated for the decomposition process observed in the present investigation. Hence it can be concluded that diffusion of manganese atoms alone is responsible for the observed decomposition process.

5. Conclusions On the basis of the above observations, the following conclusions can be drawn. (1) Mechanical alloying of elemental aluminum and manganese powders led to the formation of a highly supersaturated solid solution of manganese in aluminum. The extent of super-

saturation was estimated to be as much as 18.5 at.% Mn. (2) The equilibrium A16Mn phase formed on elevated temperature exposure of the powder mix. The volume fraction of the AI~,Mn phase increased from about 0.06 at 6 2 3 K to 0.21 at 723K. (3) Room temperature aging of the as-milled powder mix resulted in the formation of a metastable f.c.c, phase with a = 0.4472 nm. (4) The activation energy for the decomposition was calculated to be 2.04 eV, close to the activation energy for the diffusion of manganese atoms in the aluminum matrix.

Acknowledgments This work was done while the authors held National Research Council-Air Force Materials Laboratory Research Associateships. They are also grateful to Dr. F. H. Froes for many useful discussions.

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