Scripta METALLURGICA Vol. 24, pp. 873-878, 1990 Pergamon Press plc et MATERIALIA Printed in the U.S.A. All rights reserved

MICROSTRUCTURE-PROPERTY CORRELATION IN AI/TiB2(XD*) COMPOSITES

A.K.Kuruvilla, K.S.Prasad, V.V.Bhanuprasad and Y.R.Mahajan Defence Metallurgical Research Laboratory

Hyderabad - 500 258, INDIA

(Received October 16, 1989) (Revised February 22, 1990)

Introduction

Aluminium alloys discontinuously reinforced with ceramic particulates, whiskers or fibres are currently being developed for various high performance applications. SiC appears to be the most popular choice by virtue of its good compatibil ity with aluminium and influence on the strength and stiffness of the composite. However, other materials like B.C(1), TIC(2), TiB2(3,4) and AI20 ~ (5) have been examined for potential reinforcement applications in aluminium m~trix composites. In a previous study (6), a comparison of the tensile properties and the moduli of AI/B4C, AI/TiC, AI/SiC and AI/TiB 2 composites showed that TiC was the most potent reinforcement. Recently Roebuck and Forno (4) have systematically carried out a comparative study of 2014-AI alloy reinforced with either SiC or TiB 9 particulates. It was found that TiB~ is a more effective reinforcement as ~ompared to SiC as far as the srengthening effect is concerned.

The size of the reinforcement particles would have an influence on the properties of the composites. It is expected that a reduction in the particle size of the reinforcement would lead to improvements in strength, assuming all other things (shape, chemistry and distribution) are equal. The matrix reinforcement bond also plays a significant role in influencing the properties of the composite. A strong mechanical bond without the presence of a chemical reaction product would be expected to show the best properties in composites.

In light of the forgoing discussion, the possibil ity of forming TiB_ z in-situ in an aluminium matrix by XD process (7) appears to be very attractlve, as it opens up a new avenue for producing metal matrix composites. In this process, Ti and B are elementall~ blended with A1 and the resultant blend is compacted and heated to about 800 C. In the presence of a liquid phase, the Ti and B react exothermically to produce submicron size TiB 9 dispersoids. These dispersoids would have a considerable effect on the strength of the composite. The matrix-reinforcement interface could also be sound considering that the reinforcement is formed in-situ in the matrix (8,9).

The present study compared two types of AI/TiB 2 composites. In one case, the AI/TiB 9 composite was prepared by blending TiB 2 particulates with AI, while in the ot~er case, Ti and B were added in the elemental form to A1 and subsequently processed so as to form AI/TiB~ composite. Such a comparison is unavailable in the literature; it is also timely since this process appears to be attractive and clear understanding of the process and nature of the interface is lacking.

Experimental Procedure

The composites were fabricated using powder metallurgical techniques. The atomised aluminium powder (0.023% Fe, 0.15% Si) was supplied by the Metal Powder Co., Madurai, India, and had an average particle size of 45um. In one case, the

* Propr ,etory process of Mart in Mar ie t ta Laborator ies

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Copyright (c) 1990 Pergamon Press plc

87~ MICROSTRUCTURE CORRELATION Vol. 24, No. 5

A1 powder was blended with TiB9 particulates (99.5% pure and average size of 4 um from CERAC) to form A1 / 20 v/o TiB 9 composite. The uniform dispersion of TiB 9 par t icu la tes in A1 matrix was achieved by following the wet blending route usiHg polar solvent (i0). In the other case, Ti powder (ii00 ppm 09, 200 ppm N~ from Micron Metals, Inc.) with an average size of 36 um was bl~nded with B powder (99.9% pure from CERAC) having an average size of 1 um, and this mixture was further blended with the A1 powder. In elemental blending, 17 wt% of Ti and 8 wt% of B powders were taken so as to form 20 volume percent of TiB- in the composite if all the Ti reacted with B. Those two powder blinds were cold-isostatica~ly pressed at a pressure of 200 MPa and subsequently degassed in vacuum at 450 C for lh. The AI-Ti-B compact was heated to 800 C in an argon atmosphere and held at that temperature for 15 mins. The AI/TiB 2 compact thus obtained (henceforth referred to as AI/TiB 2 (XD)) was porous in nature. One sample 8f pure A1 powder was also cold-isostatically pressed, degassed in vacuum at 450 C and sintered at 600C for lh. This was done for the purpose of comparison with the composites.

Both the compacts were rolled into rods of 6mm dia at a temperature of 600C. Rolling of these compacts into round sections was accomplished by a special canning process which helped to consolidate the material and prevent breakage while rolling. The processing sequence for the AI/TiB 2 composite is schematically shown in Fig.l(a), and that for AI/TiB 2 (XD) is shown in Fig.l(b).

Samples were taken from the composite rods and polished for metallographic examination. The microstructure was observed both under an opt ica l microscope and SEM. Thin foils for TEM were prepared by a twin jet polishing technique, using an electrolyte of methanol, butanol, and perchloride acid. The foils were observed in a Philips EM430T electron microscope at 300 kV. X-ray diffractometer scan was made on AI/TiB 2 (XD) specimen using Cu Ku radiation.

Elastic moduli of these samples were measured on machined rods of 5mm dia and 100mm length using a resonance technique (by analysing the vibrational behaviour of the sample following an impulse excitation). Tensile test samples having a gauge diameter of 4 mm and a gauge length of 25mm were machined from the rods and tested at a strain rate of 6.6 x i0 /s. The fracture surfaces of the tensile tested samples were examined in the SEM.

Results and Discussion

The microstructure of the AI/TiB9 composite is shown in Fig.2. The TiB 2 particulates appear well dispersed in he matrix, though they are irregular in shape. In comparison, the AI/TiB 2 (XD) composite reveals a different microstructure as shown in Fig.3. THere are some coarse particles which have been found to be an intermetall ic phase of unknown composition A1 Ti. Apart from these coarse precipitates, micrometer -sized TiB 2 precipitates ar~ also present, and adhere well to the matrix.

Table 1 shows the tensile properties of the two composites. The tensile property values of pure A1 are also shown for comparison. The elastic modulus of the AI/TiB 9 composite (96 GPa) is substantially higher than that of pure A1 (70 GPa) althSugh the value is much lower than the value of 143 GPa predicted by application of the rule of mixture. However, the AI/TiB 9 (XD) composite shows a value of 131 GPa, which is substantially high and is almost double the modulus of pure AI. Considering the strength values, the AI/TiB 2 composite displays a yield strength of 121 MPa compared to the 64 MPa of pure A1 and a UTS of 166 MPa compared to the 90 MPa of a pure A1. The AI/TiB 2 (XD) composite displays dramatic improvements in strength. The YS and UTS of this composite are 235 MPa and 334 MPa, respectively. These strength values are signif icantly higher than that of a conventionally processed A1/TiB 2 composite. The major contribution to stengthening is most likely due to the flne dispersion of the high-modulus TiB 2 particles which are well bonded to the matrix(8).

The XD route followed in the present case probably involves the solution of

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Ti and B in liquid A1, and the solute assisted reaction leading to the in-situ formation of submicrometer size TiB9 dispersion(7). Eventhough the exact mechanism is not clear at this time, ~uch a process as described above would result in the formation of TiB^ in the liquid phase, thus forming an excellent matrix/reinforcement bond. Th~ experimental results of XRD (Fig.4) show the presence of TiB~ along with A1 peak in AI/TiBo (XD) composite. The calculated 'd' values of TiB 9 are very close to the theor~tica! values. This observation is further confirmea with TEM of the same composite specimen. It may be noted that the X-ray diffraction peaks corresponding to the A1 Ti phase could not be observed as its volume fraction was below the deection limits o f x-ray analysis.

The microstructure of the AI/TiB~ (XD) composite is shown in Fig.5. The TiB~ particles a.e present in ~he =orm of hexagonal and block-llke crystals (Fig.5=7 and the corresponding SADP is shown in Figure 5b. These TiB~ particles are of sizes ranging from a sub-micrometer to a micrometer level ~nd are distributed uniformly in the matrix. A high dislocation density is observed at the particle matrix interface (Fig.6). The presence of a high dislocation density observed at the interface is probably due to (i) semi-coherent AI-TiB 2 interface produced by in-situ formation of TiB~ crystals and (ii) thermal shock associated with the exothermic reaction. The presence of extremely fine size TiB~ particles in the matrix along with the dense interfacial dislocation sub-strucEure results in the increased strength level in the AI/TiB~ (XD) composite. It may be noted that the conventionally produced composite ~ontains coarser and incoherent TiB 2 particles, which are probably responsible for the lower strength.

The fracture surface of the tensile samples show the presence of micron-sized TiB^ in the case of the AI/TiB 2 (XD), as seen in Fig.7. Essentially the fracture surface displays dimples which are small as compared to that of the conventionally processed AI/TiBo composite (Fig.8). Since the size of the dimples indicate matrix controlled fracture defined by the interparticle spacing, the smaller dimples suggest the fineness and closenes of the dispersion of TiB 2 in the AI/TiB 2 (XD) composite.

Conclusion

The XD process(7) shows great promise for the fabrication of metal matrix composites for high performance applications. The AI-20 v/o TiB9 (XD) composite shows a modulus of 131 GPa, which is higher than that of conventionally produced composite. This process also results in MMCs with superior strength properties, as compared to conventionally processed 5~Cs. These results are attributed primarily to the presence of extremely fine size TiB~ particles, excellent matrix reinforcement bonding, and dense interfacial disl~cation sub-structure.

Acknowledgements

The authors thank Dr.P.Rama Rao, Director, Defence Metallurgical Research Laboratory, for his encouragement and permission to publish this paper.

References

l.H.K.Slater, P.M.Coon, Proc. 4th Int. A1 Extrusion Tech. Seminar, Chicago, (1988), 2, p.525.

2.Baturinskaya et al., Izv. Akad. Nauk SSSR Met., (1983), 3, p.166. 3.J.W.Mccoy. C.Jones and F.E.Wawner, SAMPE Quarterly, (1988), 19 (2), p.37. 4.B.Roebuck and A.EJ.Forno, "TiB 9 and SiC Particulate Reinforced Aluminium

Alloy Metal Matrix Composites", i~ Modern Developments in Powder Metallurgy, Voi.20, p.451, Proc. 1988 International Powder Metallurgy Conf., June 5-10, 1988.

5.David M.Schuster, Michael D.Skibo and Will iam R.Hoover, Light Metal Age, (1989), Feb., Voi.47, p.15.

6.A.K.Kuruvilla, V.V.Bhanuprasad, K.S.Prasad and Y.R.Mahajan (To be published in Bulletin of Materials Science, 1989).

7.L.Christodoulou, D.C.Nagle and Brupbacher (1986) Int. Pat. No. WO86/06366.

876 MICROSTRUCTURE CORRELATION VoL. 24, No. S

7.L.Christodoulou, D.C.Nagle and Brupbacher (1986) Int. Pat. No. WO86/06366. 8.T,M.F.Ronald, Advanced Materials and Processes, (1989), 135 (5), p.29. 9.AR.C.Westwood, Met. Trans. A, (1988), !gA, p.749.

10."Method of Making SiC Whisker Sheet Composites", U.S.Patent No.4 634 608, January 1987.

Table - 1

Material E YS UTS % E1 Hardness (GPa) (MPa) (MPa) VHN

Pure A1 70 64 90 21 37 A1/20 V/O TiB2 96 121 166 16 85 AI/20 V/O TiB2(XD) 131 235 334 7 110

AI Ti B 2 \ /

I WET BL END ING

COLD ISOSTATIC ] PRESS ING

L J DEGASSING }

450C/lh .

SINT~RING I I 6oo'c/,. ]

1 [ CANNING I

SEC;ION ROLLING

,, I I COMPOSITE L ROD

At Ti.B \ /

I DRY BLENDING

I . COLD I505TATIC

PRESSING

l DEGASSING 450"C/1h

I REACTION SINTERING 800"Cllh

I CANNING I

SECTION ROLLING

I I COMPOSITE

ROD

- =-

(a) (b)

Fig.2: SEM micrograph showing TiB 2

particulates in conventionally

processed A!/TiB 2 composite

%

o=

m

#

Fig. I: Schematic representa- tion of processing sequence

(a) AI/TiB 2 conventional

processing

(b) AI/TiB 2 XO process

Fig.3:SF24 micrograph showing

blocky AlxTi intermeta!!ics and

T i~ particles in AI/TiB 2 (XD) composite

Voi. 24, No. 5 NICROSTRUCTURE CORRELATION 877

! E I

O

11

i ~L ___jL_" j k_. . . . . . .

20

Fig.4: X-ray diffractogram showing TiB 2 peaks

Fig.6: TEM mlcrograph showing interracial dislocation sub- structure

Fig.5: (a) TEJ~ micrograph showing the TiB 2 crysta l~ in AI/TiB2(~) composite

(b) SADP of TiB9 - [O00I] Zone Axis

878 MICROSTRUCTURE CORRELATION Vol. 24, No. 5

Fig.7: SEM fractograph of A!/TiB 2 (XD) composite

Fig.8: S~ fractograph of AI/Ti8 2 composite