Materials Science and Engineering, A 153 (1992) 691-695 691

Synthesis of A1/A13Ti two-phase alloys by mechanical alloying

S. Srinivasan, S. R. C h e n and R. B. Schwarz Center for Materials Science, Los Alamos National Laboratory, Los Alamos, NM 87545 (USA)

Abstract

We have mechanically alloyed mixtures of elemental powders to prepare fine-grained two-phase AI/AI3Ti powders at the compositions A1-20at.%Ti and Al-10at.%Ti. Hexane was used to prevent agglomeration of the powder during mechanical alloying. Carbon from the decomposition of the hexane was incorporated in the powder and reacted with titanium to form a fine dispersion of carbides in the final hot-pressed compact. We consolidated the mechanically alloyed powders by hot pressing. Yield strength and ductility were measured in compression. At 25 °C, the compressive yield strengths were 1.25 and 0.6 GPa for the AI-20at.%Ti and Al-10at.%Ti alloys respectively. The ductility of the Al-10at.%Ti alloy exceeded 20% for 25 < T< 500 °C.

1. Introduction

Aluminum-rich intermetallics such as A13X (X = Ti, Zr, and Nb) are being studied because they possess an important combination of properties: low density, good oxidation resistance, and strength and modulus reten- tion at high temperatures [1-3]. These intermetallics crystallize in ordered low-symmetry tetragonal D022 or D023 structures which give the alloys high strength and modulus. However, they are extremely brittle near ambient temperature. Recent research has focussed on developing ordered alloys that retain the attractive high temperature properties whilst improving the low tem- perature ductility. Researchers are following a two- pronged approach to solve this problem: (1) solute additions in the atomic percent range (macroalloying) aimed at stabilizing the alloy in the cubic L12 struc- ture-related to the D022 and D023 structures [4], and (2) very dilute solute additions (microalloying) aimed at decreasing the twinning and antiphase boundary (APB) energy in the D022 and D023 phases to improve the low temperature ductility of these phases [2, 5].

The first approach has spawned a large number of studies, particularly in the low density AI3Ti system. It has been found that additions of copper [6-8], iron [9], manganese [6, 10, 11], chromium [11] and nickel [12, 13] stabilize the L12 phase in A13Ti. Similarly, iron, chromium, nickel and copper have been added to AI3Zr to increase the stability of the L 12 phase [14, 15]. Most of these L 12 alloys have been tested in compres- sion, showing only a few percent ductility at room

temperature. Liu [1] has recently shown that some A13X-based L12 alloys that show ductility in compres- sion are generally brittle in tension. Nevertheless, the research community still considers approach (1) as most promising for the eventual achievement of low density, high strength, high temperature alloys. The second approach for improving the ductility of the D022 and D023 phases has been more eagerly pursued by theoreticians. Fu [16], Nicholson et al. [5] and Yamaguchi [2] have studied the role of APB and twin- ning energies in the plastic behavior of AI3Ti (D022 structure). In their view, twinning and slip would be facilitated by decreasing these energies (possibly through microalloying), leading to a more ductile behavior.

We are investigating fine-grained two-phase metal-intermetallic alloys to develop high strength alloys useful to intermediate temperatures, The ap- proach is similar to that used in the design of oxide dispersion strengthened nickel-based superalloys: we hope to achieve strength using an intermetallic phase, achieve ductility by having fine grains and a second ductile phase, and prevent grain growth by having a fine dispersion of stable precipitates. The precipitates should also improve the high temperature creep resis- tance. There is interest in two-phase AI/A13X (X = Ti, Zr, and Hf) alloys, also known as super "alumalloys", which are expected to be useful to 75% of the absolute melting point of aluminum [ 17, 18].

In this study we report the synthesis and properties of two-phase A1/AIsTi alloys with different aluminum

0921 - 5093/92/S 5.00 © 1992 - Elsevier Sequoia. All rights reserved

692 S. Srinivasan et al. / Synthesis of Al/Al~Ti two-phase alloys

contents. The alloy powders were prepared by mechanically alloying mixtures of elemental powders and were consolidated by hot pressing.

2. Experimental details

The alloy powders were prepared by mechanical alloying (MA). MA is a high energy ball-milling process originally developed for preparing oxide dispersion strengthened nickel-based alloy powders [19]. MA produces fine-grained powders which are homogene- ous in composition. In addition, MA often leads to metastable crystalline or amorphous phases [20].

The MA was done in a Spex 8000 laboratory mill using a hardened steel vial and lid, and 50 hardened steel balls of diameter 6.35 mm. The starting materials were a mixture of 99.99% pure aluminum (from Cerac, Milwaukee, WI) and 99% pure titanium (from Johnson Matthey, Ward Hill, MA). The vial and the lid were sealed with an elastomer O-ring. With MA there is always the danger that the powder may become con- taminated by iron eroded from the steel milling media. By weighing the cleaned milling media after a large number of MA runs we deduce that for a 5 g powder charge and 20 h of MA (typical of our runs) the iron contamination was less than 0.001 g per run.

Approximately 6 ml of hexane (C6H14) was used as a dispersant to prevent the agglomeration of powder on the walls of the container and on the balls. Using hexane, the powder recovery was approximately 90%. Following the milling, the hexane was removed by evaporating it in a partial vacuum. Typically, for 20 h of MA, a 5 g powder charge retains between 0.10 and 0.20 g of hexane, which decomposes during MA and is incorporated into the powder mostly as homogene- ously distributed elemental hydrogen and carbon [15, 21]. X-ray diffraction (XRD) also shows that during MA some of the hydrogen combines with titanium to form Til l 2. Both forms of hydrogen are easily removed during a vacuum anneal near 500 °C. During subse- quent hot pressing at temperatures above 500 °C, the elemental carbon, approximately 2 wt.%, combines with titanium to form a fine dispersion of TiC.

The handling of the elemental and alloyed powders (loading of the milling vials, recovery after mechanical alloying and loading of the graphite dies) was done inside a glove-box filled with argon continuously re- circulated through a getter. The gas had less than 0.1 p.p.m. 02 and 0.1 p.p.m. H20 (by volume). We have not measured the final oxygen content in the hot-pressed compacts but we expect it to be low, close to that of the starting commercial powders.

The structures of the powders were determined by XRD using Mo Ka or Cu Ka radiation. Further experi-

mental details have been published elsewhere [15, 18, 21, 22].

Two-phase AI/A13Ti alloys were prepared by hot pressing the mechanically alloyed powders in a Grafoil-lined graphite die. The two compositions of Al /Alf f i reported here are (1) AI-20at.%Ti (hence- forth Als0Ti20 ) and (2) Al-10at.%Ti (henceforth A190Til0 ). The hot-pressing methods for the two alloys were slightly different.

The mechanically alloyed AlsoTi20 powder was annealed in an ultrahigh vacuum furnace for 40 rain at 480 °C to remove the hydrogen introduced during MA. The hot pressing was done in a flowing argon atmo- sphere at 830 °C, using a pressure of 12 MPa. Part of the excess aluminum squeezed out of the die as the pressing temperature was well above the melting point of aluminum. The volume fraction of aluminum re- tained in the compact was deduced from the measured heat of fusion of the aluminum phase (Fig. 1 ). From the amount of hexane retained in the powder, and assum- ing that all the carbon forms TiC, we calculated the carbide content in the product. Thus, the final com- position of the AlsoTi20 alloy was Al f f i -1 lwt.°/oA1 - 9wt.%TiC. This alloy had a density of 3.25 g cm- 3.

The A190Til0 powder was consolidated in a vacuum chamber at 650 °C under a hot-pressing pressure of 34 MPa. The hydrogen was removed from the powder during hot pressing prior to the application of pressure by holding it at 500 °C until the vacuum was better than 50 #m. In this alloy, all the aluminum in the mechani- cally alloyed powder was retained in the final compact. The final composition of the Alg0Til0 alloy was A1-28wt.% A13Ti-8wt.%TiC and had a density of 2.92 g cm -3.

The as-pressed samples were disks 31.7 mm in diameter and 6.5 mm in height. Compression speci-

60 I I

O

A

v

~= 45

I

35

,30 , , , I ~ , , I ,

200 400 600 800 TEMPERATURE (*C)

Fig. 1. Heat evolved during the first heating of MA AlsoTi2o powder measured by DSC.

S. Srinivasan et aL / Synthesis of Al/Al~Ti two-phase alloys 693

mens, 5 mm in diameter and 6 mm in height, were cut from the disks by electro-discharge machining. Com- pression tests were done in an MTS-800 system at a strain rate of 0.0002 s -~. The specimens were placed between polished platens of Udimet-700, using Teflon tape as lubricant for tests up to 400 °C and fine glass powder for tests above 400°C. The yield stress in compression was determined from the intersection of tangents drawn to the elastic and the initial plastic portions of the stress-strain curves. Samples for trans- mission electron microscopy (TEM) were prepared by diamond polishing to a thickness of 50 ~m followed by ion milling with the specimens cooled to 80 K.

3. Results

We have measured the lattice parameters of meta- stable and stable phases of AI3Ti [18], and we have used these values to calculate the lattice misfit between the A13Ti and A1 phases. Comparisons with previously measured and calculated values show close agree- ments.

Figure 2 is a scanning electron micrograph of an etched surface of hot-pressed Als~Ti20 alloy. Narrow channels of aluminum surround AI~Ti regions. These channels are preferentially oriented in planes perpen- dicular to the hot-pressing axis. This morphology resulted from the extrusion of the excess aluminum during the hot pressing. XRD patterns of the compacts show that the A13Ti phase has the stable D022 struc- ture. The transmission electron micrograph in Fig. 3

In the as-mechanically alloyed condition, the A13Ti powders have the cubic L12 structure [18, 21, 22]. This structure is metastable and upon heating, the alloys transform to the stable D022 structure. We have observed [18, 22], however, that before the L12-A13Ti transforms to the stable D022 phase, it transforms to a second metastable D023 phase, intermediate in free energy to the L12 and D022 phases.

Figure 1 shows the heat evolved during the first heating of Al80Ti20 powder which was mechanically alloyed for 20 h. These measurements were taken by differential thermal analysis (DTA) in a Perkin-Elmer (model 1700), at a heating rate of 2 0 K s -1. The analyzer was operated in the differential scanning calorimetry (DSC) mode using the a-fl transformation of manganese or the fusion of aluminum as tempera- ture and enthalpy calibration standards. The reactions giving rise to the peaks in the DSC traces were iden- tified by taking XRD patterns of powder that was heated to temperatures just below and above these peaks. The first exothermic peak A, between approxi- mately 400 and 550 °C, is generated by the LI 2 ~D023 transformation of the A13Ti phase. Knowing that this phase is 80 wt.% of the compact, we deduce that the L12~D023 transformation enthalpy is close to 7 kJ (mol atom)- ~. Endothermic peak B is generated by the melting of the excess aluminum. The peak correspond- ing to the D023 ~D022 transformation of the AI3Ti phase is not seen, probably being hidden by peak B. However, measurements for single-phase mechanically alloyed AI3Ti give a D023--'D022 transformation enthalpy of 0.9 kJ (mol atom)-1. Freeman and co- workers [23] have calculated from first principles the formation enthalpies for the D019 , L12 and D022 phases of AI3Ti. From these values we can estimate the L12 --'D022 transformation enthalpy to be 5 _+ 1 kJ (mol atom) -l.

Fig. 2. Scanning electron micrograph of hot-pressed Als0Ti2o alloy, polished and etched to reveal aluminum inclusions.

Fig. 3. Transmission electron micrograph of hot-pressed AlsllTi21 ~ alloy.

694 S. Srinivasan et aL / Synthesis of Al/Al3Ti two-phase alloys

shows an average grain size of approximately 0.5/~m for both the aluminum and AI3Ti grains. The TiC precipitates, 30 nm in size, were found dispersed on the surface of the A13Ti grains. The surface of the aluminum grains were relatively free of carbide pre- cipitates.

The results of the compression tests on the Al80Ti20 samples are shown in Fig. 4. The room temperature yield strength (open circles) is in excess of 1 GPa, decreasing with increasing temperature. There is a stronger decrease as the test temperature approaches 500 °C, most likely due to the softening of the pure aluminum phase. The plastic strain to fracture (open triangles) is zero at room temperature but increases with test temperature. At 500 °C, the sample did not fracture even at 80% plastic strain. The solid symbols are the compression test results of Yamaguchi e t al. [2] for coarse-grained AI3Ti (1 mm grain size). Both the yield strength and the ductility of the present alloy exceed those of the coarse-grained single-phase A13Ti alloy. We attribute the difference in properties to the difference in grain size (0.5 vs. 1000/~m), the presence of a ductile aluminum phase and the fine dispersion of carbides (30 rim) in our material.

The mechanical properties of the Al80Ti20 alloy, having only 11 wt.% AI because of the way it was pre- pared, indicate that a larger volume fraction of alu- minum is needed to improve the room temperature ductility. For this reason we mechanically alloyed a powder mixture at the composition Al90Ti~0. Further- more, to retain all the excess aluminum, we chose a consolidation temperature of 650 °C. Figure 5 shows a scanning electron micrograph of the polished surface of the Alg0Ti]0 sample in the backscattered electron imaging mode. X-ray energy dispersive analysis shows

1 .5 , , i , , i , , , i , , , i , , 10

that the darker phase, Seen in the micrograph as a minority phase, is pure aluminum. The lighter regions are the A13Ti phase. Preliminary TEM observations show AI3Ti grain sizes between 0.15 and 0.20 pro. Closer observation of the scanning and transmission electron micrographs shows a finer distribution of aluminum between AI3Ti regions. This is expected because for this alloy the pure aluminum is the major phase, close to 70% by volume.

Figure 6 shows the temperature dependence of the compressive yield strength (open circles) in Alg0Tiz0. For all temperatures, the stress-strain curves were smooth, showing little, if any, work hardening. The tests were stopped at 20% plastic strain, indicating that the fracture strain exceeds this value. The yield strength at room temperature is 600 MPa, decreasing

Fig. 5. Scanning electron micrograph of the polished surface of hot-pressed Al9oTi m alloy in the backscattered electron imaging mode. The darker regions correspond to aluminum and lighter regions to AI3Ti.

0 " - " - 0 ~ I 8 o ~ x

O 3 - J

0 3 - -

n , -

4 ° jO .5 _ _z

I P - - • I

o o = ' - - , ~ , , , t i 1 0

400 600 800 1000 TEMPERATURE (K)

Fig. 4. Compressive yield strength (o) and plastic strain to fracture (zx) of A180Ti2o alloy as a function of temperature. The closed symbols are data for large-grain single-phase AI3Ti from ref. 2.

700

6OO

o

500

400 ( . 9

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300

a

-- 200 W >-

I oo

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0 Al-16.5wt.~.Ti-l.57*C, MA

1 O0 200 300 400 500 TEMPERATURE (°C)

Fig. 6. Compressive yield strength of AIg0Ti,, alloy as a function of temperature.

S. Srinivasan et al. / Synthesis of Al/Al3Ti two-phase alloys" 695

to 100 MPa at 500 °C. These yield strengths are well above those observed for commercial A1-Cu (2024) and AI -Mg (7075) alloys. The compressive yield strengths of our alloy are close to the tensile yield strengths of two of the advanced aerospace aluminum alloys recently developed at Allied-Signal [17, 24] by rapid solidification. We will have a more meaningful comparison after testing our alloys in tension.

The present results on A180Ti~0 and Alg0Ti~0 show that the increase in the aluminum content leads to an increase in the room temperature ductility. Optimizing the aluminum content may enable us to develop alloys of high strength whilst retaining an acceptable degree of ductility.

4. Discussion

Yamaguchi et al. [2, 25] identified the (111)[112] twinning to be the deformation mode in AI3Ti at tem- peratures below 620 °C. Because there are only four independent twinning systems, plastic deformation generates incompatibility-related stress concentrations. We believe that the presence of aluminum surrounding the AI3Ti grains provides a sink for the relief of these stresses leading to a pseudo-augmentation of the defor- mation. Yamaguchi et al. [25] also suggest keeping the twin area fine to increase the ductility of AI3Ti at low temperatures. The fine microstructural scale of our mechanically alloyed material should provide such a condition. These observations could explain why the ductility in our alloys is significantly larger than that of coarse-grain single-phase AI3Ti. T E M analysis of the deformed samples is now underway to confirm these explanations.

There is considerable interest in developing low density, high strength, high oxidation resistant materi- als useful to temperatures close to 425 °C. Several studies [17, 18, 26] have identified the A1/AI3Ti and A1/A13Zr systems as having the highest potential for developing these super "alumalloys" because: (1) the aluminum and AI3X phases have' a low lattice mis- match [18, 26], and (2) titanium and zirconium have a relatively low diffusivity in aluminum. The results of this study are encouraging and support this view. The advantageous combination of strength and ductility in the present alloys derives from a cumulative effect of the intrinsic strength of the AI3X phase, the micro- structural refinement and the fine carbide dispersion.

Further work is necessary to understand the deforma- tion mechanisms of these alloys.

Acknowledgment

This work was supported by the US Department of Energy, Office of Basic Energy Sciences.

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