SINTERING BEHAVIOR OF MECHANICALLY ALLOYED Ti-48Al-2Nb ALUMINIDES
© 2014 D. D. Mishra*, V. Agarwala and R. C. Agarwala
Dept. of Metallurgical and Materials Eng., Indian Institute of Technology Roorkee, Roorkee-247667, India
*e-mail: debeshmaterials@gmail.com Received February 14, 2013
Abstract—Ti—Al intermetallics have been produced using mechanical alloying technique. A composition of Ti—48Al—2Nb at % powders was mechanically alloyed for various durations of 20, 40, 60, 80 and 100 h. At the early stages of milling, a Ti (Al) solid solution is formed, on further milling the formation of amorphous phase occurs. Traces of TiAl and Ti3Al were formed with major Ti and Al phases after milling at 40 h and beyond. When further milled, phases of intermetallic compounds like TiAl and Ti3Al were formed after 80 hours of milling and they also found in 100 h milled powders. The powders milled for different durations were sintered at 785°C in vacuum. The mechanically alloyed powders as well as the sintered compacts were characterized by XRD, FESEM and DTA to determine the phases, crystallite size, microstructures and the influence of sintering over mechanical alloying.
DOI: 10.7868/S0040364414010153
INTRODUCTION
Ti-Al based intermetallics are recognized as the materials having desirably low density, high strength to weight ratio, high stiffness and high temperature oxidation resistance for potential use in high structural materials for automotive and aerospace applications [1—6]. The room temperature ductility is the only obstacle which inhibits the forming of big structural parts, so different methods are applied to improve that: grain refinement, chemical treatment and thermo mechanical heat treatment [7, 8]. Mechanical alloying (MA) is one of the advantageous processes which gives rise to a ductility, strength and nanostructure of the intermetallics with respect to other powder metallurgical processes, which results in the formation of fine nanoscale grains and homogeneous particles distribution. During mechanical alloying fracturing and re-welding of the powder particles occur with incipient fusion, which creates non-equilibrium amorphous solid solutions [9—12]. Fadeeva et al [13] and Bhat-tachraya et al [14] reported the formation of metasta-ble fcc phase during annealing of Ti (Al) solid solution. Sintering of these aluminides includes two processes: reactive and non-reactive. Reactive sintering takes place due to chemical reactions occurring at particular interface resulting in the production of new material, while the later one is the consolidation and recrystalli-zation of mechanically alloyed solid solution.
The aim of this study is to determine the effect of milling on the morphological evolution of intermetal-lic compounds formed during MA and sintering pro-
cesses. Another aim is to investigate the effect of MA on the densification of the intermetallics.
MATERIALS AND METHODS
Ti (40-44 ^m, 99.9% purity), Al (40-44 ^m, 99.9% purity) and Nb (40-44 |im, 99.9% purity) powders were mixed to give the composition Ti-48-Al-2Nb (at %) which was then charged into a hardened steel vial with hardened steel balls under a wet toluene media i.e. the balls and charge are totally submerged in the toluene. The charge to ball weight ratio (CBR) was 1 : 5. The milling was performed in a (Retsch PM 400/2) ball mill for periods varying from 0 to 100 h at a milling speed of 300 rpm and vial rotation speed of 600 rpm (Table 1).
Process control agents, such as toluene, were used to prevent agglomeration of elemental powders. For the purposes of reactive and non-reactive sintering it was performed on two different series of powders: (1) reactive sintering process done on blended elemental powders, and (2) non-reactive sintering process done on 100 h milled powders. The powders were sintered under ultra high pure argon environment for 2 h at 450°C and 800°C .The sintered samples were 5 mm in thickness and 16.4 mm in diameter. The milled and sintered samples were characterized by means ofX-ray diffraction (XRD) using a D8 BRUKER AXS diffrac-tometer with Cu Ka, operating at 40 kV and 30 mA. Scherrer's formula was used to determine the crystallite size with the help of X-ray line broadening [15]:
d = 0.9Vpcos0. (1)
Table 1. Milling parameters
Planetary Ball mill Details (Retsch PM 400/2) Milling Parameters
Milling Balls — Hardened steel balls Milling Media — Toluene
Milling Jars — Hardened steel jars Charge to Ball ratio — 1 : 5
Jar capacity — 500 mL Milling speed — 300 rpm
Vial Speed — 600 rpm
PCA — Toluene Total time of milling — 100 h
Weight of initial charge — 30 g
Where p = (p
-й f
PM is the full width at half maxima (FWHM), P7 is the instrumental broadening correction factor, 9 is the angle of the maxima, and X is the Cu K wavelength (0.15406 nm).
The thermal analysis of the MA powders of different milling times and as-received powders were done up to 1400°C in a PERKIN ELMER Pyris Diamond TG/DTA. DTA (Differential Thermal Analysis) measurements were carried out in argon atmosphere to avoid oxidation. The argon flow rate was kept 200 mL/min and the heating rate was kept 10°C/min. QUANTA FEI-200 Field emission scanning electron microscopy (FESEM) was used to characterize the morphology of the milled powder and the surfaces of the sintered samples. Energy-dispersive analysis of X-Rays (EDAX) coupled with FESEM was used for the semi-quantitative investigation of the microstructure of the sintered samples. The densification of the sin-
tered samples was calculated using the Archimedes principle. According to that principle when a body is immersed into fluid there is buoyant force in upward direction on the body which has the same weight of that of the displaced fluid
Fb = mLg,
FB is buoyant force and mL is mass of the liquid displaced. The dimensions of the samples were measured by the help of vernier calipers and the mass of them was measured. Then the density of the samples is calculated in air by using the equation p = m/V.
Then the sample is submerged into water and the weight of the sample is taken when submerged.
After that the actual density of the sample is calculated by using the equation
Pa = WJWa — Ww)),
Intensity 300
250
200 -
150 -
100 -
50
20
30
40
50
60
70 80
29, degrees
0
Fig. 1. XRD patterns of milled powders for different times: 1 - Ti, 2 - Al, 3 - TiAl, 4 - Ti3Al, 5 - 20 h milled, 6 - 40, 7 - 60, 8 — 80, 9 - 100.
Al
Crystallite Size, nm 180 160 140 120 100 80 60 40 20 0
Al, Ti(Al) solid solution (TiAl, Ti3Al) Ti(Al) solid solution (TiAl, Ti3Al) TiAl
TiAl
ml
20 40 60 80 100 Time of Milling, h
Fig. 2. Variations of crystallite sizes of the MA Ti-48Al-2Nb at different milling times.
DTA
2 0 -2 -4 -6 -8 10 -12 14 16 18
цУ (а)
200 400 600 800 1000 1200 1400 Temperature, °C
Temperature, °C
800 700 600 500 400 300 200
20 40 60 80 100 Time of Milling, h
Fig. 3. Thermal analysis of as-received Ti-48Al-2Nb powder mixture (a) and (b) the starting (3), peak (2) and the finishing (1) temperatures of the exothermic peak in the thermal analysis data of powders of different milling times.
0
0
where pa is actual density, Ww is weight in water and Wa is weight in air.
RESULTS AND DISCUSSION
MA process
XRD analysis. The XRD Fig. 1 depicts the evolutions of transformation occurred during milling of Ti-48Al-2Nb. The XRD pattern of as-received powder is almost similar to that of 40 h milled powders. The peak broadening and lowering in the intensity was observed with increase in extent of milling, due to the decrease in crystallite size occurring from 91-177 nm after 20 h to about 12-18 nm after 100 h of milling. Distortion of lattice parameter of Ti results in shifting of the main reflexion of Ti peaks occurring towards higher angles. From the Ti-Al phase diagram it can be seen that Al has a better solubility in Ti, while the reverse is limited for different temperatures. The amorphization occurs
because of the free energy of intermetallics is higher than that of amorphous phase [16]. So as per the expectation, Al highly diffused into Ti in the process of MA up to a certain extent and beyond which an amorphous phase is formed. At this stage peak broadening leads to the overlapping of peaks, as a result of which the peaks are not easily distinguished. The amorphiza-tion of the powders and the lowering of the peak intensities can be attributed towards the external energy supplied to the system by mechanical alloying [17].
According to Fig. 1, after 40 h of MA, the peaks are broadened between angles of 34—42°; which indicates the amorphization of powders. Interestingly the free energy curve of the amorphous phase in the Ti—Al system is actually lower than that of the solid solution and intermetallics. Hence, the amorphous phase is formed first, then further milling leads to the formation of supersaturated solid solution hcp-Ti (Al) due to the crystallization of an amorphous phase or disordered Ti3Al,
Fig. 4. FESEM micrographs of as-received and milled powders for different milling times: (a) as-received powder, (b) 10 h MA powder, (c) 20, (d) 40, (e) 60, (f) 80 and (g) 100.
because these two phases present the same unit cell [14]. The diffraction pattern after 80 h of MA shows the formation of the TiAl intermetallic compound. After 100 h of MA, both TiAl and Ti3Al peaks are present in the diffraction pattern. The enthalpy values for the formation of TiAl and Ti3Al are —75 and —73 kJ/mol respectively [18—20] which is the main cause of formation of intermetallics and that further heat treatment is required for formation of intermetallics after MA [17]. In addition the local temperature rises during MA which helps the nucleation of the intermetallic compound. The MA of the powder mixtures results in the lowering of the crystallite size due to the external force applied by ball milling and the plastic deformation occurring. The crystallite size of the as-received powder of
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