In situ fabrication of Al3Ti particle reinforced aluminium alloy metal–matrix composites

Materials Science and Engineering A364 (2004) 339–345

In situ fabrication of Al3Ti particle reinforced aluminium alloy metal–matrix composites Xiaoming Wang a,∗ , Animesh Jha b , Rik Brydson b a

Department of Metals and Materials Engineering, McGill University, 3610 University Street, Montreal, Que., Canada H3A 2B2 b Department of Materials, University of Leeds, Leeds LS2 9JT, UK Received 15 May 2003; received in revised form 21 August 2003

Abstract Titanium tri-aluminide (Al3 Ti) particles were dispersed homogeneously into a castable aluminium alloy matrix by the aluminothermic reduction of hexafluorotitanate (K2 TiF6 ) under different conditions. Al3 Ti particles in different morphologies and sizes were produced by changing the fabrication conditions, such as composition of the flux, the temperature and holding time. The coarsening and growth of the Al3 Ti particulates are principally affected by other elements present in the flux during fabrication. The effects of the temperature and holding time, alloying elements and the composition of flux on the dispersion of the reinforcement were examined by using SEM and X-ray diffraction techniques. The observed results are explained in terms of the different growth behaviour of the Al3 Ti particles under different conditions. The dispersion of the Al3 Ti particles and the Al/Al3 Ti interfacial bonding are explained by the solidification of aluminium. © 2003 Elsevier B.V. All rights reserved. Keywords: MMCs; Solidification; Aluminium; Composite; Al3 Ti; Al–Ti

1. Introduction Metal–matrix composites (MMCs) have been attractive amongst other types of structural engineering materials for more than 40 years. MMCs normally comprise a metallic matrix and a ceramic reinforcement in the form of particles, whiskers or fibres. These materials have the combined properties of their constituents, therefore, can be tailored to offer improved properties to meet different engineering requirements. Two of the important properties of MMCs are their improved mechanical and wear-resistant properties [1]. Aluminium alloys reinforced with ceramic particulates, such as Al2 O3 , SiC and TiB2 , etc., have long been investigated because these composites exhibit higher specific properties compared with the traditional monolithic materials. MMCs have mainly been produced via powder metallurgical methods due to the poor wettability of the ceramic reinforcement by liquid aluminium [2]. These techniques are expensive compared with liquid-phase processing methods. The high costs arising either from processing or the com∗ Corresponding author. Tel.: +1-514-398-4755x0086; fax: +1-514-398-4492. E-mail address: [email protected] (X. Wang).

0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2003.08.049

ponents restrict the application of MMCs to certain areas where cost is not the prime consideration, such as aerospace and military components [1]. The interface between the matrix and the reinforcement plays an important role in determination of the properties of the MMCs [1]. The interface acts as a bridge for load transfer from the matrix to the reinforcement. Strong interfacial bonding between metal–matrix and reinforcement is desirable for high performance MMCs. Therefore, it is important to eliminate the formation of a brittle layer at the matrix/reinforcement interface. The brittle layer at the interface is usually formed by the oxidation of the reinforcements or by reactions between matrix and reinforcement during the fabrication steps or under high temperature service conditions [3,4]. For example, Al4 C3 forms in Si-free Al-alloys whilst Al2 MgO4 forms in Mg-containing alumina reinforced aluminium composites. On the other hand TiB2 particle reinforced Al MMCs offer a high degree of promise due to their good interfacial properties [5–7]. However, to disperse exogenously formed TiB2 particles into an Al-melt is difficult. Recently, in situ formed TiB2 particle reinforced Al-alloy MMCs have been successfully fabricated by using a fluoride flux assisted method, which originally arose from the grain refinement practice in aluminium industries [8–10].

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The nucleation of ␣-Al from TiB2 is believed to proceed indirectly through an Al3 Ti layer [11]. This provides the inspiration for using Al3 Ti as the reinforcement to fabricate Al3 Ti reinforced aluminium MMCs. As candidate reinforcement, in situ Al3 Ti particles offer lighter weight and stronger direct bonding with the Al-alloy matrix compared with in situ TiB2 particulates. The density of Al3 Ti is about 3.3 g/cm3 , which is comparable with that SiC of 3.27 g/cm3 . Al3 Ti being an ordered phase also demonstrates high strength and elastic modulus (∼220 GPa) [12]. The intermetallic phase also has an excellent resistance to oxidation and corrosion in fluoride atmosphere well above the melting point of Al. The intermetallic structure can be produced via casting technique, as has been cited in the Al–Ti–B master alloy literature [13]. The high resistance to oxidation of Al3 Ti offers a strong capability of this MMC to resist corrosion in application, especially, at the reinforcement/matrix interface. The reinforcing phase of Al3 Ti is produced via the reaction of molten aluminium and hexafluorotitanate (K2 TiF6 ), as shown in reaction (1). Therefore, the in situ fabrication technique for MMCs can be easily adapted with the existing equipment that has been used in the aluminium industry for conventional casting. K2 TiF6 +

13 3 Al

= Al3 Ti + 23 (3KF · AlF3 ) + 23 AlF3

(1)

In this paper, the effects of the fabrication parameters, including the effects of alloying elements, on the formation and dispersion of the in situ formed Al3 Ti particles are discussed and the mechanism of the microstructural evolution is explained.

2. Experiments The in situ experiments were designed on the basis of reaction (1) for producing the Al3 Ti intermetallic phase in the Al-alloy matrix. In each experiment, 10–100 g of commercial purity aluminium (or aluminium alloys) were placed inside a high-alumina crucible, which was covered by a lid and melted in air using a resistively heated muffle furnace. When the molten aluminium reached the processing temperature (950 ◦ C), K2 TiF6 flux was added slowly onto the surface of the melt. After the addition of flux, the melt was maintained at the same temperature for 30 min, during which time, the

formation, dispersion and growth of the intermetallic particles were progressing. To investigate the effects of temperature and other fluxes on the growth of the Al3 Ti particles, temperature were raised to 1050 ◦ C and other fluxes were then added into the crucibles. The mixtures were then kept at the same temperature for a fixed time period varying from 1 to 3.5 h. Al-alloy containing intermetallic particulates were poured into steel moulds and cooled to room temperature in air. Experimental details are listed in Table 1. The MMC ingots were sectioned for microstructural analysis. The samples for SEM analysis were polished to a 1 ␮m finish and examined using a Camscan Series 4 SEM operated at 20 kV. The elemental distributions of titanium and aluminium in the cast ingots were investigated using energy dispersive X-ray (EDX) analysis. X-ray diffraction (XRD) analyses were performed on discs of cast MMC samples and powders of fluxes using a Philips APD 1700 Diffractometer.

3. Results and discussions 3.1. Reaction between the flux and aluminium alloy Fig. 1a and b show the typical microstructure of the in situ Al–Al3 Ti MMCs produced via the reaction of K2 TiF6 with aluminium (sample MT001). Fig. 1a is a low magnification SEM back scattered electron micrograph, which shows the overall uniform dispersion of blocky Al3 Ti particulates in the aluminium matrix. High magnification micrographs, as shown in Fig. 1b, reveal that Al3 Ti particulates are within the ␣-Al grains. The presence of Al3 Ti particles is confirmed by XRD analysis, Fig. 2. The homogeneous dispersion of the in situ Al3 Ti particulates can be explained by the crystal structures of ␣-Al and Al3 Ti and the Al–Ti phase diagram (Fig. 3). The Al–Ti phase diagram shows that there is a peritectic reaction between the aluminium melt and the Al3 Ti phase at the aluminium end of the phase diagram. The solidification of aluminium in the presence of Al3 Ti occurs by forming a layer of solid aluminium on the surface of the Al3 Ti particulates, and subsequent diffusion of titanium into the ␣-Al layer [14]. The in situ formed Al3 Ti particulates are therefore in the centres of ␣-Al grains. The diffusion of titanium also forms a diffusion Al/Al3 Ti bond, and offers improved interfacial bonding strength compared to the mechanical bonding in most

Table 1 Experimental details for producing in situ Al–Al3 Ti MMCs Sample

MT001 MT003 M1005 M1006 a

Composition (g)

Additives (g)

Holding

Al

K2 TiF6

MgF2

KBF4

LiBF4

Li2 TiF6

Li2 ZrF6

Time (h)

Temperature (◦ C)

20 10 100 25

10.4 1.3 26.0a 13.0

– – 2.3 2.3

– 3.1 26.0 –

– – – 13.0

– – 5.2 –

– – 1.3 –

– – 3.5 3.0

– – 1050 1050

Containing 20 g K2 TiF6 and 6 g K2 ZrF6 .

X. Wang et al. / Materials Science and Engineering A364 (2004) 339–345

341

Fig. 3. Aluminium-rich corner of the Al–Ti phase diagram [14,18].

Fig. 1. The microstructures of in situ Al–Al3 Ti MMCs: (a) low magnification SEM image showing homogeneous dispersion of the in situ Al3 Ti particulates and (b) high magnification SEM image showing in situ Al3 Ti particulates in the ␣-Al grain centers, temperature 950 ◦ C, time 30 min (sample MT001).

casting MMCs. The pronounced tendency of Al3 Ti particulates in nucleating ␣-Al grains arises due to the superlattice structure of Al3 Ti. The tetragonal crystal structure of Al3 Ti (a = 0.3848 nm and c = 0.8596 nm) is comparable to that of ␣-Al (a = 0.4049 nm). As a result, the interfacial strains in both the a and c directions of Al3 Ti, εaa = (1−aAl3 Ti /aAl )× 100 = +4.96% and εcc = (1 − cAl3 Ti /2aAl ) × 100 = −6.15%, are small. Such calculations point out that the ␣-Al lattice will form a coherent boundary with Al3 Ti. This crystallographic consideration is confirmed by the HREM analysis of the Al3 Ti/Al/TiB2 system reported by Schumacher et al. [15]. Because of the coherent interface between ␣-Al

300 - Al - Al3Ti

250

Counts

200

150

100

50

0 10

20

30

40 50 Two theta (degree)

60

70

80

90

Fig. 2. XRD pattern of Al–Ti alloy produced by reaction of K2 TiF6 with aluminium showing diffraction peaks for an Al3 Ti phase (sample MT001).

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grains spontaneously during solidification, which results in a homogeneous dispersion of the reinforcing particulates. 3.2. Effects of alloying elements on Al3 Ti phase morphology

Fig. 4. Flaky Al3 Ti phase produced in the presence of KBF4 , showing a large aspect ratio compared to sample MT001 (Fig. 5).

and Al3 Ti, it is expected that an Al–Al3 Ti composite may exhibit comparable strength to those of Al–TiB2 and Al–SiC cast composites. SEM analysis also reveals the absence of reaction products and impurities in the Al/Al3 Ti interfacial region. Therefore the detrimental effects of segregation of impurities and reaction products in the matrix/reinforcement interface, which were discussed in different MMCs by Rohatgi and Asthana [4], are also diminished. Al3 Ti particulates in two different morphologies have been produced in the in situ Al–Al3 Ti MMCs, both nominally equiaxed (blocky) and plate or needle-like (flaky) as shown in Figs. 1 and 4, respectively. The blocky particulates have a small aspect ratio, whilst the aspect ratio of the flaky Al3 Ti particulates is large. The variation in the morphologies of Al3 Ti was explained on the basis of growth behaviour under different conditions in the Al–Ti–B aluminium grain refiners [13]. The flaky Al3 Ti forms by growing along its 1 1 0 direction of dendrite arms at high temperature and moderate cooling rate, whilst the blocky particulates grow in both 1 1 0 and 0 0 1 directions under supersaturation of titanium in the Al–Ti–B melt. The crystal structure of Al3 Ti is shown in Fig. 6 [13]. When there are no other alloying elements and the flux is a single composition of K2 TiF6 , the in situ Al3 Ti particulates are mainly blocky in shape. The reduction of K2 TiF6 yields a high supersaturation of titanium in the aluminium melt, especially in the Al/flux interface, and results in the formation of blocky Al3 Ti particulates. The blocky Al3 Ti particulates are the most effective nucleation sites for ␣-Al grains amongst their three different morphologies, flaky, blocky and petal-like, because blocky Al3 Ti crystals exhibit (0 1 1) planes, which offer low crystal mismatch with the (0 1 2) planes in ␣-Al [13,16]. This gives the highest probability for the blocky Al3 Ti particulates to be in the ␣-Al grain centres. The uniform particulate size of the blocky Al3 Ti particulates means that they will all have similar ability in nucleating ␣-Al grains. Therefore, the Al3 Ti particulates nucleate ␣-Al

3.2.1. Effects of boron containing flux Alloying elements that modify the Al/Al3 Ti interfacial energy, such as magnesium, lithium [17] and boron, play important roles in determining the morphology of the in situ Al3 Ti particulates and consequently their dispersion. In situ Al3 Ti with different morphologies were produced by adding magnesium, lithium or boron in the form of fluoride salts: LiBF4 , MgF2 and KBF4 . Fig. 4 is an SEM micrograph of a typical sample produced in the presence of KBF4 . It shows the existence of flaky Al3 Ti particulates with a very large aspect ratio. This indicates that KBF4 depresses the growth of Al3 Ti along the 0 0 1 direction and speeds up the growth along the 1 1 0

direction. When there is no boron reduced from KBF4 in the melt, Al3 Ti grows without any preferred direction and appears more equiaxed in shape. The effect of boron on the morphology of the Al3 Ti particulates in aluminium matrix is more complex than that predicted by thermodynamic calculations, which suggest that Al3 Ti decomposes to form TiB2 [18]. The depressed growth of Al3 Ti in the 0 0 1 direction can be attributed to the attack of boron on the (0 0 1) planes, which is supported by the decomposition of the Al3 Ti phase along the elongated surfaces in the presence of Zr, seen in Fig. 5. It suggests that the Al3 Ti particulates were attacked by the boron-rich melt and form TiB2 particulates. From the crystal structure of Al3 Ti, shown in Fig. 6, Ti atoms mainly occupy the (0 0 1) planes. Therefore, in the presence of boron these planes will be attacked first because of the large negative free energy change for the formation of TiB2 [18]. In the presence of TiB2 , Al3 Ti nucleates on the surface of the TiB2 , and causes the transfer of titanium from Al3 Ti particulates to TiB2 particulates.

Fig. 5. Decomposition of Al3 Ti particles from their (0 0 1) plane.

X. Wang et al. / Materials Science and Engineering A364 (2004) 339–345

-Ti;

- Al

343

Fig. 8. SEM micrograph showing the coarsening of Al3 Ti particulates in the presence of LiBF4 and MgF2 (sample M1006).

Fig. 6. Crystal structure of Al3 Ti [13].

3.2.2. Effect of surface modifiers in the cryolite melt It is well known in the literature of Al-alloy casting that the addition of Li and Mg in the alloy or in the flux during melting can substantially modify the surface tension of the liquid alloy [17]. Lithium at about 0.4 wt.% in liquid aluminium reduces the surface tension by one-third at its melting point, which implies that the presence of lithium will also alter the interfacial energy of the solid–liquid interface, resulting into a major change in the nucleation and growth rates of the solid phase. The effect of Mg in liquid aluminium is comparable with lithium. Mg, however, does not depress the surface tension as much as Li in liquid aluminium. The effect of Mg in the presence of a fluoride flux

is to substantially reduce the surface modification tendency of Mg by forming a stable fluoride complex, KMgF3 , which locks Mg atoms. The formation of KMgF3 is evident from the X-ray diffraction pattern of solidified flux, shown in Fig. 7. The morphology of Al3 Ti, which was produced in the presence of LiBF4 and MgF2 , is shown in Fig. 8. It manifests three effects in stages: (a) first an increase in the nucleation rate of Al3 Ti phase is observed due to the presence of either Li atoms dissolved in the alloy phase or due to the presence of LiBF4 in the flux, which provides the source of Li+ ions at the flux–melt interface; (b) the bi-axial growth of Al3 Ti along the 0 0 1 and 1 1 0 directions take place;

600 - KMg F 3

- Al

500

Counts

400

300

200

100

0 25

35

45 Two theta (degree)

55

65

Fig. 7. An XRD pattern of the flux produced by reactions of K2 TiF6 and KBF4 with Al–8 wt.% Mg alloy showing the formation of KMgF3 , temperature 950 ◦ C, time 30 min.

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B KBF4

A

<110> direction (a) Original Al 3Ti particles

(b) Elongated Al3Ti

0 0 1 direction. Boron attacks the Al3 Ti particulates on the (0 0 1) plane and this results in the high interfacial energy of the plane, thus makes the mass transport on this plane easier. When there is MgF2 together with LiBF4 , the surface condition of the (0 0 1) and (0 1 1) planes of particulate A is modified and the growth of new layers occurs on these planes by consuming particulate B, as shown in Fig. 10.

Fig. 9. Illustration of the effect of boron on the growth-morphology of in situ Al3 Ti particulates.

4. Conclusions (c) finally the columnar growth of Al3 Ti with concomitant formation of (Al,Ti)B2 solid solution phase in the intergranular regions of ␣-Al. The columnar growth appears to be sustained via the consumption of smaller blocky Al3 Ti particulates. Unlike the conditions shown in Fig. 5, there is no excess boron present in the liquid, which can promote the attack on the needle and lath-like Al3 Ti, since boron atoms are used up in the latter stages of reactions to form (Al,Ti)B2 . From Fig. 8, it is evident that the larger the reinforcing particulates the stronger their ability to nucleate ␣-Al grains, results in the smaller particulates being pushed into ␣-Al grain boundaries during solidification. This unique microstructure is emphasised when LiBF4 and MgF2 were used to modify the morphology of the reinforcing particulates. As shown in Fig. 8, the presence of LiBF4 and MgF2 fluxes improves the growth of the Al3 Ti particulates dramatically. Compared with sample MT001 (Fig. 1), LiBF4 and MgF2 make Al3 Ti particulates grow hundred times larger in size. When MgF2 addition was made together with LiBF4 , both (0 0 1) and (0 1 1) planes of Al3 Ti are modified to be suitable for the growth of new layers of Al3 Ti phase, therefore blocky Al3 Ti particulates form and grow larger. TiB2 may also form on the (0 0 1) plane of the Al3 Ti particulates and nucleate a new layer of Al3 Ti. The smaller particulates in the ␣-Al grain boundaries are evidence for these boride (AlB2 ) particulates, which form later than the Al3 Ti particulates. The mechanisms by which the alloying elements, boron, lithium and magnesium, change the morphology of in situ formed Al3 Ti particulates are proposed from the observed microstructures and are shown schematically in Figs. 9 and 10. In Fig. 9 particulate A grows in the 1 1 0 direction by consuming particulate B and mass transfers along its own

B

LiBF4 & MgF2 A (a) Original particulates

(b) Coarse Al3Ti particulate

Fig. 10. Illustration of the effects of LiBF4 and MgF2 on the growthmorphology of Al3 Ti particulates.

In situ Al3 Ti particle reinforced aluminium metal matrix composites have been produced successfully. The reinforcing particles dispersed in the aluminium matrix homogeneously, due to their strong ability to nucleate ␣-Al grains. This also makes the interfacial bonding between these two phases strong. Particle size and morphologies are largely affected by the processing conditions, such as temperature, holding time and the composition of the flux. High temperatures and longer holding times result in the coarsening of the particles. However, the growth behaviour changes in the presence of surface-active elements such as magnesium, lithium and boron, which lower the interfacial energy in particular crystal directions of the particles and leads to a change in the particle morphology.

Acknowledgements The authors acknowledge the CVCP U.K. and the University of Leeds for granting an ORS Award and a TetleyLupton studentship to X. Wang for his study in Leeds.

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