TMS Outstanding Student
Structural Evolution in Mechanically Alloyed Al-Fe Powder Mixtures
Debkumar Mukhopadhyay, author
CONTENTS
The structural evolution in mechanically alloyed binary aluminum-iron powder mixtures containing 1, 4, 7.3, 10.7, and 25 at. pct Fe has been investigated using x-ray diffraction and electron microscopic techniques. The constitution (number and identity of phases present), microstructure (crystal size, particle size) and transformation behavior of the powders on annealing have been investigated. The solid solubility of Fe in Al has been extended up to at least 4.5 at. pct. compared to the equilibrium value of 0.025 at. pct Fe at room temperature. A fully amorphous phase plus solid solution in the Al-10.7 at. pct Fe alloy; agreeing well with the predictions made using the semi-empirical Miedema model.
Nanometer sized grains were observed in the as-milled crystalline powders in all compositions. Heat treatment of the MA powders containing the supersaturated solid solution or the amorphous phase resulted in the formation of the Al3Fe intermetallic in all but the Al-25 at. pct Fe powders. In the Al-25 at. pct Fe powder, formation of nanocrystalline Al3Fe2 was observed directly by milling. Electron microscope studies of the shock-consolidated mechanically alloyed Al-10.7 and 25 at.pct Fe powders indicated the nanometer-sized grains were retained after compaction.
A number of "far from equilibrium" synthesis techniques have been developed during the past few years to produce materials with improved properties; these include rapid solidification (RS) processing, vapor deposition, and mechanical alloying (MA). Even though the principles utilized in these various processing techniques are quite different, the main objective of all of them is the same - microstructural and constitutional flexibility control thereby allowing enhanced physical and mechanical properties. Amongst these, MA has been shown to be a powerful technique to produce non-equilibrium phases, in addition to its traditional application to produce oxide-dispersion strengthened (ODS) nickel- and iron-base alloys.[1.2]
The presence of these non-equilibrium phases allows greater flexibility and control of the final equilibrium microstructure, thereby producing materials with properties superior to those obtained by more conventional ingot metallurgy and casting techniques. The non-equilibrium phases synthesized include supersaturated solid solutions, metastable crystalline or quasicrystalline intermediate phases, and amorphous phases.[2,3]
Very often, and depending on the conditions prevailing, these processing techniques produce microstructures with highly refined grain size (down to nanometer levels), precipitate size, or segregation patterns.[4,5] Mechanical alloying is a solid-state powder processing technique which involves repeated welding, fracturing and rewelding of powder particles in a dry, high energy ball mill.[1,2]
An inherent advantage of the MA process is that it is a solid state process, thus allowing alloying of constituent elements of significally different melting points or those which are immiscible in the solid state; which is often difficult or impossible by conventional melting techniques.
Alloying of Al with Fe increases the high temperature strength due to a dispersion of second phase particles. This effect can be further enhanced by increasing the volume fraction of the second phase obtained though increased solid solubility extension of Fe in Al by techniques such as RS or MA. The equilibrium solid solubility of Fe in Al at room temperature is only 0.0025 at. pct*, Fig 1.[5] Intermetallic compounds have very high melting points and therefore can be used for high temperature structural applications but they are brittle at room temperature. However, it has been shown that nanostructured materials have improved ductility over their coarse-grained counterparts.[7]
Amorphous phases possess an interesting combination of properties and provide several potential advantages. the amorphous structure can also be used as a precursor to synthesize nanostructured material and the microstructure can be controlled as desired. Thus, the intent of this paper is to report results of our extensive investigations on MA of Al-Fe alloys containing from 1 to 25 pct Fe and critically analyze the results on the formation of intermetallic compounds and metastable phases (supersaturated solid solutions and amorphous phases) in this alloy system.
Mechanical alloying was earlier applied to the Al-Fe system to produce supersaturated solid solutions and amorphous phases in different alloy compositions.[8-13] These composition ranges are different depending on the severity of the milling (type of mill) (Table 1).
Synthesis of intermetallics was not generally achieved directly by mechanical alloying; but, a subsequent annealing is required. However, there are some reports of synthesis of intermetallic compounds by milling in the Al-Fe system. [13,14]
* All compositions in this paper are expressed in at. pct unless otherwise mentioned.
Comp. range of Maximum solid Amorphous phase of powder solubility formation range Type of Mill Ref. (Pct Fe) (Pct Fe) (Pct Fe)4-42 1 17-35 Horizontal 8,9 3.5 3.5 - - 10 10-60 - 20-50 Spex 11 20 - 20 Spex 12 12.5-25 - 20 Planetary 13
Nominally pure Al (>99.4% and -100 mesh) and Fe (>99.4% and -325 mesh) powders were used as starting powders to make up mixtures containing 1, 4, 7.3, 10.7 and 25 pct Fe, the last composition corresponding to stoichiometric Al3Fe intermetallic. MA was carried out at room temperature in a SPEX 8000 Mixer mill. The grinding medium was 3/16" diameter hardened 52100 steel balls. For each run approximately 10 g powder mixture and 100 g steel balls were used and milled using forced air cooling.
Small quantities of the powder were withdrawn from the vials at regular time intervals inside an argon-filled glove box to follow the progress of alloying via X-ray diffraction (XRD), transmission electron microscopy (TEM) and scanning electron microscopy (SEM) techniques. The milled powders were heat treated by sealing them in evacuated glass tubes to evaluate the stability of the phases formed. The mechanically alloyed powders were compacted by shock consolidation method and the compacted specimens were also characterized by XRD and TEM.
General Observations:
In all the five alloys investigated (Al-1, 4, 7.3, 10.7 and 25 pct Fe), a decrease in the particle size was observed with increasing milling time. Fig. 2 shows SEM micrographs of the Al-4 pct Fe powder in as-blended condition and after milling for 10h and 20h, indicating that the particle size decreases with increasing milling time. A similar situation was obtained in all the other compositions.
At a high magnification (Fig.3) it can be seen that the observed particles are composite in nature and are made up of agglomerated nanometer-sized particles; the particles are less than 100mm in size in the Al-10.7 pct Fe powder MA up to 30h.
The XRD patterns in each case showed that the Al(200) and Al (220) peaks overlap with Fe (110) and Fe (200) peaks, respectively; and only the Al (111) and Al (311) peaks are seen clearly and separately. Increasing milling time caused the peaks to broaden and decrease their intensities due to both the refinement of crystal size and introduction of strain during milling. In powders with high Fe contents, the Al (111) peak is found to have a lower intensity than the Al (200) peak, whereas at lower Fe contents, the Al (111) peak is more intense, as expected.
The crystal size was calculated for each condition from the broadening of X-ray peaks using the Scherrer formula [15] and subtracting the strain and instrumental broadening effects. Nanometer-sized crystals are formed in all the alloys after long MA times (Fig.4). Increasing the ball-to-powder weight ratio(BPR) resulted in a faster rate of decrease of particle size; a lower BPR decreased the rate of reduction.
The results of investigation regarding the formation and stability of metastable phases in the powders with different Fe contents will now be presented.
A. Solid solution formation:
Fig.5 shows the XRD patterns of the Al-7.3 pct Fe powder as a function of milling time. In the starting powder mixture all the expected peaks from both Al and Fe can be seen with the Fe (110) and Fe (200) peak overlapping with the Al(200) and Al(220) peaks, respectively. With increasing milling time the Al peaks shifted to higher angles suggesting alloying of Fe with Al leading to formation of a solid solution.
The lattice parameter of the Al solid solution, calculated from the position of the Al (311) peak, decreased in 20h from the initial value of 0.40497 nm in the as-blended state ot about 0.4022 nm for the Al-4 pct Fe powder and to 0.4015 nm for the Al-7.3 pct Fe powder and thereafter remained constant indicating that a steady state condition has been reached . The steady state lattice parameters in different powder compositions were used to evaluate the solid solubility in each case using the master plot of lattice parameter vs. Fe content (Fig. 6). From this it is clear that a maximum of 4.5 pct Fe could be dissolved in Al in the solid state by MA.
B. Synthesis of Intermetallics:
Fig.7 shows the XRD patterns of the as-milled Al-25 pct Fe powder. After 15h of milling partial formation of a compound was observed directly by milling.[14] The JCPDS data suggested that this compound can be identified as Al5Fe2 with an orthorhombic structure and a = 0.767 nm, b = 0.64 nm, and c = 0.42 nm. 100% formation of Al5Fe2 compound was observed after 30h of milling; the peaks are relatively broad due to the fine crystal size and the presence of strain in the powder. On annealing the powder milled for 30h at 625 degrees Celsius for 324h, sharp x-ray diffraction lines resulted (Fig.8), including a number of them with low intensities which did not show up in the as-milled powder. However, there is no difference either in the structure or the lattice parameters of this phase between the as-milled and annealed conditions.
In the other compositions, however, formation of the intermetallic phases was not observed directly by milling. Fig.9, the XRD pattern of the Al-10.7 pct powder MA for 50h and heat treated at 625 degrees Celsius for 3h shows the presence of both Al and Al3Fe. Similar heat treatment for the 4 pct Fe powder Ma for 20h showed formation of Al3Fe compound. The relative proportion of the Al and Al3Fe phases obtained after heat treatment in these compositions matches well with that calculated from the equilibrium design.
C. Amorphous Phase Formation:
As can be seen from Fig.7, an amorphous phase was obtained after milling the Al-25 pct Fe powder for 50h. Electron diffraction techniques confirm the formation of the amorphous phase in the powder, Fig.10. However, only partial amorphization was observed in the Al-10.7 pct Fe powder after 50h of milling. Amorphization was not observed in Al-1, 4 and 7.3 pct Fe compositions.
D. Consolidation:
The Al-10.7 pt and 25 pct Fe powders MA for 65h and 15h, respectively, were consolidated by shock companion techniques [16] to form pellets. Near 100% density was achieved in both the specimens. X-ray diffraction patterns of the consolidated specimens indicate the presence of Al3Fe and Al3Fe2 compound, respectively, in the 10.7 and 25 pct Fe powders suggesting that the solid solution phase has transformed to Al3Fe in the powder with 10.7 pct Fe and to Al5Fe2 in the 25 pct Fe during shock consolidation. Transmission electron micrographs of the consolidated specimens (Fig.11) confirmed the retention of 25-30nm grains in both the compositions after the shock consolidation.
A summary of the results of MA in different compositions was shown in Fig.12.
Mechanical alloying of Al-Fe alloys has been shown to result in particle/grain size refinement, extended solid solubility limit of Fe in Al, intermetallic compound and amorphous phase formation.
General Observations:
A comparison of the XRD patterns (Fig.5 and Fig.7) of Al-7.3 pct Fe and Al-25 pct Fe powders, shows that in the case of Al-7.3 pct Fe powder, the Al (111) peak is more intense than the Al (200) Fe (110) peak even after 20h, and Al (220) Fe(200) and Al (311) peaks are still present. But in the case of Al-10.7 pct and higher Fe contents the Al(111) peak has a much lower intensity than the Al(200)Fe(110) peak after long milling times.
After continued milling the Al(311) peak has completely disappeared and the Al(220)Fe(200) peak is very broad and has low intensity. Similar results were reported earlier in MA Al-Cr [17,18] and Al-Mo [19,20] systems. All the solute elements used in these investigations have a bcc structure. Materials with a bcc structure are generally brittle in comparison with the fcc aluminum. During milling, the brittle Fe powder gets fragmented more easily than Al, and gets coated on to the surface of the ductile aluminum. Since X-ray diffraction occurs from a finite depth from the surface it is then likely that we observed the diffracted intensity mostly from the bcc Fe phase and only a little from the Al phase.
In the case of lower concentration of Fe, this did not happen because the amount of the brittle phase was not sufficient to coat the aluminum surface to the same extent as at higher Fe contents. A very similar result was also observed during milling of brittle TiH2 and ductile Al [21] electron microscopic confirmation for the above proposal was obtained. It would be useful to test this hypothesis using a ductile metal to be alloyed with aluminum.
Solid Solubility Extension:
The solid solubility limit of Fe in Al achieved in the present investigation is higher than that observed by Huang and co-workers.[8,9]
The higher energy spex mill used by us is probably responsible for this higher amount of supersaturation. Polkin et al.[10] observed a solid solubility of 3.5 pct Fe in Al by MA, but due to lack of any detailed information (e.g. type of mill, ball:powder ratio) in their paper further discussion is not possible. A number of investigators reported solid solubility extension of Fe in Al by RS and these results are summarized in Table II and compared with those obtained by MA.
It can be seen from the table that the solid solubility extension obtained in the Al-Fe system by MA is at least as high as the maximum extension observed using RS techniques. In RS techniques, large extension of solid solubility limits are possible by quenching the alloy melts to below the To temperatures (temperatures at which the solid and liquid phases have the same free energy at a given composition) [28] and a metastable equilibrium is reached between the supersaturated solid solution and the intermetallic (if present in the system).
However, in MA technique, the repeated welding and fracturing of the powder particles results in a very intimate mixture of the two constituent powder particles with a concomitant increase in the defect concentration and strain in the powder mixture. Continued milling refines the grain size further and an even greater number of defects are introduced. Since the diffusion distances are reduced, the grain boundary area is increased and the defect density is high, diffusion rates are enhanced and true alloying occurs between the two constituents. However, in MA a metastable equilibrium is reached between the supersaturated solid solution and the amorphous phase and this limits the solid solubility.[29]
Equilibrium solid Extension Extension
solubility by RS Ref. by MA Ref.
(Pct Fe) (Pct Fe) (Pct Fe)
0.025 0.082 22 1 8,9
3 23 3.5 10
3.5 24 4.5 Present Study
3.7 25
4 26
4.4 27
Synthesis of Intermetallics:
In the Al-25 pct Fe powder MA for 15h produces the Al5Fe2 intermetallic compound directly by milling. Since the starting powders were of the overall Al3Fe composition, EDS analysis was conducted on the he MA powders to check whether the composition corresponds to stoichiometric Al3Fe composition. The analysis indicates that the powder contains 73.5 pct Al and 26.5 pct Fe. This change in composition can be either due to loss of Al due to sticking of the powder on the he container wall and to the grinding balls and/or due to gain of Fe from the he steel balls and container wall.
Since the A5Fe2 intermetallic forms under equilibrium conditions in the composition range of about 27 to 29 pct Fe[6] it is not surprising that formation of homogeneous Al3Fe2 was observed in the present investigation at an Fe content of 26.5 pct. Similarly, Huang [8] also observed the formation of the Al5Fe2 intermetallic in the nominal Al-24.4 pct Fe powder MA for 180h, but only after heat treatment at 500 degrees Celsius.
To determine whether the Al3Fe intermetallic can be synthesized directly by MA after compensating for the loss of Al, an extra amount of 1.5 pct Al was added to the nominal Al-25 pct Fe powder mix prior to MA. However, the Al3Fe intermetallic phase did not form; instead, a solid solution of Fe in Al formed at early milling times an an amorphous phase formed after 50h. Although the enthalpy of formation of Al3Fe (28.1 kJ/g-atom) is very close to that of Al5Fe2 (28.26 kJ/g-atom), Al3Fe could not be directly synthesized probably due to the complex crystal (monoclinic) structure and large size of the unit cell.
The Miedema model predicts that an amorphous phase should form in the composition range 25 to 60 pct Fe in the Al-Fe system. Accordingly an amorphous was observed in the Al3Fe composition in the present work, which is one of the boundaries of the composition range predicted by Miedema model [30] and by other investigations. [8,11]
The intermetallic Al3Fe compound could, however, be synthesized in the present investigation after heat treating the MA powders. However, Morris and Morris [13] reported the formation of Al3Fe compound directly by ball milling using a Fritsch mini planetary ball mill at a ball:powder ration 6:1 in the Al-20 pct Fe and Al-25 pct Fe compositions. They observed formation of an amorphous phase which, on continued milling, transformed to the Al3Fe intermetallic in the Al-20 pct Fe powder.
But in Al-25 pct Fe powder the Al3Fe compound was directly observed without any intermediate amorphous phase. Continued milling by Morris and Morris [13] beyond the time for the formation of the Al3Fe compound in Al-25 pct Fe powder may have resulted in the amorphous phase formation; although the lower milling intensity (and hence reduced energy in the system) used by these workers may have precluded formation of the amorphous phase, which normally requires severe deformation and accumulation of sufficient defect concentration.
Amorphous Phase Formation:
Miedema's semi-empirical model [30] has been used to predict the glass forming range in the Al-Fe system. The enthalpy vs. composition diagram calculated using the Miedema model is shown in Fig.13. The common tangent construction suggests that a fully amorphous structure should form between 25 and 60 pct Fe. At compositions, x < 15 pct Fe and x> 90 pct Fe, only a solid solution is expected to be stable. In the intermediate ranges 15 pct Fe < x < 25 pct Fe and 60 pct Fe < x < 90 pct Fe both the amorphous phase and solid solution are predicted to coexist.
Table III presents the composition ranges where a fully amorphous phase can be obtained both according to the theory and experiments. This Table shows that the predicted and observed composition ranges are very close in the work of Dong et al. [11], but not for the data reported by Huang [8]. Wang et al. [12] and Morris and Morris [13] observed full amorphization in the Al-20 pct Fe composition, which is outside the composition range predicted by the Miedema model.
In the present investigation MA of Al-25 pct Fe showed full amorphization which agrees closely with the prediction. In the Al-10.7 pct Fe composition a mixture of an amorphous phase and solid solution was observed which also matches well with the prediction of the Miedema analysis. It should be noted that experimental circumstances may also play an important role in the outcome of the investigations; the type of ball mill, the ball:powder ratio, the purity of the atmosphere during milling etc. may be important factors and it is imaginable that different investigators could obtain different results in the same alloy system depending on the parameters used. Although it is recognized that the Miedema model [30] has limitations (e.g. the model does not consider the presence intermediate phases in the system), it is still the best available model at the present time to predict the glass forming range.
Comparison of Predicted Fully Amorphous Composition Range Using Miedema
Model, with the Observed Range in the Al-Fe System
Predicted (pct Fe) Observed (pct Fe) Reference
25-60 17 to 33 8
20 to 50 11
20* 12
20* 13
25** Present Investigation
Based on the results obtained on the mechanical alloying of Al-Fe alloys with 1, 4, 7.3, 10.7 and 25 pct Fe, the following conclusions can be drawn:
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Figure 1: Al-Fe Equilibrium Diagram
Figure 2: Scanning electron micrographs of the Al-4 pct Fe powder (a) as-blended (b) milled for 10h and (c) milled for 20h showing a decrease in the particle size with milling time.
Figure 3: Scanning electron micrograph of the Al-10.7 pct Fe powder MA for 30h, showing hte presence of nanometer sized particles.
Figure 4: Plot showing the variation of crystal size with milling time in AI-4, 10.7 and 25 pct Fe powders
Figure 5: X-ray diffraction patterns of Al-7.3 pct Fe powder as a function of MA time.
Figure 6: Master plot showing the variation of lattice parameter of Al-solid solutions with iron content obtained by RS and MA techniques.
Figure 7: X-ray diffraction patterns of Al-25 pct Fe powder as a function of MA time.
Figure 8: X-ray diffraction patterns of the Al-25 pct Fe powders, (a) mechanically alloyed for 30h, (b) mechanically alloyed and heat treated at 625 degrees C for 324h.
Figure 9: X-ray diffraction patterns of the Al-10.7 pct Fe powder for 50h and subsequently annealed at 625 degrees C for 3h showing the presence of Al and Al3Fe phases.
Figure 10: Electron diffraction pattern of Al-25 pct Fe powder MA for 50h showing full amorphization.
Figure 11: Transmission electron micrographs of the shock consolidated specimens of (a) Al-10.7 pct Fe powder MA for 65h and (b) Al-25 pct Fe powder MA for 15h showing the presence of 25-30 nm nanocrystals.
Figure 12: A plot showing the summary of results of MA in different powder mixtures.
Figure 13: Calculated enthalpy vs. composition diagram for the Al-Fe system using the Miedema model.
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