Synthesis and microstructural characterization of al mg alloy sic particle composite

Available online at www.sciencedirect.com

Materials Letters 62 (2008) 2623 – 2625 www.elsevier.com/locate/matlet

Synthesis and microstructural characterization of Al–Mg alloy–SiC particle composite S. Valdez a,⁎, B. Campillo a,b , R. Pérez a , L. Martínez b , A. García H b a

b

Instituto de Ciencias Físicas—Universidad Nacional Autonoma de México, Av. Universidad S/N, Col. Chamilpa, 062210, Cuernavaca, Morelos, Mexico Facultad de Química—Universidad Nacional Autónoma de México, Cd. Universitaria, 04510, México, D.F., Mexico Received 21 September 2007; accepted 3 January 2008 Available online 11 January 2008

Abstract Al–Mg-alloy was reinforced with 10 vol.% SiC particles size of 3 μm diameter by vortex technique. Stiffener distribution, particle interaction with metal-matrix and mechanical properties in as-cast condition was studied. The resulting as-cast composite structures were analyzed using Xray diffraction (XRD) and scanning electron microscopy (SEM). The AlMg–SiCp composite microstructure showed excellent SiCp distribution into AlMg matrix. In addition, no evidence of secondary chemical reactions has been observed. Hence, mechanical properties are highly sensitive to the microstructure and these are indirectly related to solidification parameters and processing conditions. Al–Mg alloy possess lightweight and excellent properties as structural materials which can be optimized with SiCp addition and a good fabrication technique. © 2008 Elsevier B.V. All rights reserved.

Keywords: AlMg alloy; Composite; Mechanical property

1. Introduction SiCp reinforced metal–matrix composites (MMCs), have been considered as excellent candidates to be applied as structural materials in the aeronautic–aerospace transport, the automotive industry, etc. [1,2]. The key to their property improvement lies in the structure [3,4], chemistry and the nature of bonding of Al–SiC interfaces. Alloying elements such as Mg, which segregate at particle–matrix interfaces, have been found to improve the wettability [3–5]. Metal–matrix composites are conventionally fabricated using different techniques such as power metallurgy, squeeze casting, and the mixing of partially solidified alloys with ceramic materials [6,7]. Stir casting tends to suffer from non-uniformity in the reinforcement distribution after solidification. Powder metallurgy is expensive [8]. An inherent difficulty encountered in the fabrication of ⁎ Corresponding author. Tel.: +52 55 56227785; fax: +52 55 56227734.

E-mail addresses: svaldez@fis.unam.mx (S. Valdez), bc@fis.unam.mx (B. Campillo), ramiroperez@fis.unam.mx (R. Pérez), lorenzo@fis.unam.mx (L. Martínez).

0167-577X/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2008.01.002

SiC–Al alloy composites is that molten Al alloys normally do not wet considerably the ceramic reinforcements. It is wellknown that the SiC reinforcements tend to react with Aluminum during processing, leading to the formation of Al4C3 and Si [9,10] at the interface. Efforts have been directed to prevent the chemical reaction at interfaces by oxidation of SiC [11,12], coating of SiC particles [13], or alloying of Al matrix with Mg or Si [14].

On the present investigation, an Al–8.7 wt.% Mg matrix alloy was reinforced with silicon carbide particles (SiCp). The SiC were added as dispersed particles by Vortex method [15,16]. The vortex method was selected in order to diminish the process cost and generating a material composite with homogenous microstructure, less casting porosity and better mechanical properties.

In addition, stringent conditions, such as high pressures, high vacuum, well-controlled atmospheres, or long fabrication times were not required to fabricate the AlMg–SiCp composite.

Composite microstructure is influenced by solidification parameters and processing conditions. Hence, tensile strength is highly sensitive to the microstructure and these are indirectly related to the preparation route, so processing parameters

2624

S. Valdez et al. / Materials Letters 62 (2008) 2623–2625

involved have a great importance. In this study, the particles distribution, phases composition and microstructure present in the aluminum–magnesium matrix are reported.

2. Experimental procedure For composite manufacture an Al–8.7 wt.% Mg (98.5% Al and 99.9% Mg of purity) was used as matrix and reinforced with 10 vol.% SiCp. The unreinforced AlMg alloy was made in an electrical resistance furnace, at 700 °C, using a flux mixture of salts (KCl + NaCl) to avoid the contact with the environment and the elements oxidation. Al–8.7 wt.% Mg matrix alloy was re-cast and preheated at 470 °C after that 10% SiC particles volume fraction was added and incorporated by vortex method at 1150 rpm for 15 min.

Microstructural characterization samples were prepared by grinding on SiC metallographic paper from 240 to 1200, polished with alumina (5.0 μm, 1.0 μm and 0.5 μm) finally, ultrasonically cleaned and attacked with Keller reagent to reveal their microstructure. The samples were examined using scanning electron microscopy (SEM) Jeol 5900 equipped with energy dispersive X-ray analyses (EDAX). X-ray diffraction was carried out using a filter of nickel with CuKα radiation in Phillips diffractometer operated to 40 kV. Tensile tests were performed using an INSTRON at a strain rate of 3 × 10− 3 s− 1 at room temperature. The tensile samples (25 × 6 × 5 mm) were fabricated following ASTM E8M-03.

3. Results and discussion The XRD pattern from the AlMg–SiCp composite manufactured by vortex method and AlMg matrix alloy (Fig. 1) shows the diffraction intensity corresponding to the α-aluminum phase, which is a solid solution rich in aluminum (α-Al) of crystalline structure FCC and the magnesium presence in low intensity. The X-ray diffraction reveals that main intensities (111) and (200) of the α-phase appear to 2θ 38.82° and 44.71° respectively. The new composite show three principal diffraction intensity to 2θ = 35.7°, 60.0° and 71.8°, which corresponds to (111), (220) and (311) of SiC-particles. In such difractograms, we can evidently deduce the crystalline phases from the master alloy and the composite material.

Fig. 1. X-ray pattern show the α-aluminum and magnesium presence on the matrix alloy Al–Mg and just SiC particles in the composite.

Table 1 Composition of AlMg–SiCp composite under study, given in at.% Phase

SSAlMg SSAlMg + SiC SiC

Elements Al

Mg

Si

C

92.72 54.56

7.28 7.12

– 17.42 52.78

– 21.20 43.23

The presence of Al2O3, Al4C3 and Si phases formed by chemical reaction between Al matrix and SiC-particles, have been reported [4]: 4Alð1Þ þ 3SiCðsÞ →Al4 C3ðsÞ þ SiðsÞ

ð1Þ

2MgðlÞ þ SiðsÞ →Mg2 SiðsÞ

ð2Þ

4Alð1Þ þ 3SiO2ðsÞ →2Al2 O3ðsÞ þ 3SiðsÞ

ð3Þ

However, in case of AlMg–SiCp manufactured with vortex method there has been no evidence of (1), (2) and (3) second reactions. Table 1 shows the phase chemical analysis obtained from the AlMg–SiCp composite. It is clear that magnesium has been entrapped into aluminum solid solution and a secondary chemical reaction product was not observed.

The explanation for non-secondary reactions, like (1), (2) and (3) may be due to the low temperature, 470 °C used during the SiCreinforcement with the vortex method. Temperatures such as 1100 °C, 800 °C, and 750 °C with 500 mmHg vacuum applied have been used in process such as power metallurgy, stirring casting technique, and infiltration technique [4,6,16]. Previous work, have been reported [4] that interfacial reaction between pure Al and SiC occurs only above 923 K (650 ◦C). Also the ternary Al–Si–C [17] phase diagram indicates that Al4C3 forms when the Al:Si ratio is below 4.5, which is not our case. In addition, the SiC incorporated speed and time do not allow diffusion of aluminium. Diffusion controls the reaction between Al and SiC. Fig. 2A show the unreinforced alloy microstructure formed by Al– Mg solid solution (SS) rich in aluminum with columnar dendritic structure. Fig. 2B shows the composite microstructure characteristic under as-cast condition with an EDAX spectrum from a spot between solid solution phase and SiC particles, evidently it suggest the absence of Al4C3 and Si phases, whose presence has been reported [4].

A homogenous microstructure can be observed avoiding particle agglomeration which has happened in advanced composite manufacture [18]. Solidification process tends to reject SiC particles close to grain boundaries, which contributes to SiC particles agglomeration. Solidification influence on particle distribution depends if SiC-particles act as solidification nuclei or SiC-particles are rejected towards interdendritic region, being in significant microstructural differences. With Vortex Method the agglomerates were not seen in the composites samples.

Fig. 2B also show the SiC particles homogenously distributed in αaluminum matrix and grain boundaries. Homogenous distribution is essential to obtain better mechanical properties [19]. Many composite materials are used by their high-specific elastic modulus, high-specific strength and good wear resistance, in addition the MMCs has been considered as potential lightweight and high-performance materials to be used in aerospacecraft, aircraft and engine parts in automobiles [20,21].

The present tensile test results (Fig. 3) clearly indicate that the strength and ductility from composite samples increase as compared to AlMg matrix unreinforced alloy, in agreement with Poddar et al. [22].

S. Valdez et al. / Materials Letters 62 (2008) 2623–2625

2625

Fig. 2. A) Al–Mg unreinforced alloy microstructure, constituted by Al–Mg solid solution rich in aluminum with columnar dendritic structure. B) AlMg–SiCp composite microstructure. Al–Mg solid solution with particles of SiC distributed uniformly and EDAX insert pattern result show the elements in the composite.

The AlMg unreinforcement specimen gives a yield stress (σ0.2) of 120 ± 5 MPa, ultimate tensile strength of 250 ± 6 MPa. Both the yield stress and tensile strength increase significantly when the AlMg alloy was reinforcement with SiC-particles. This improve tensile properties are proposed to be strongly determined by α-Al solid solution strengthening originated by SiC-particles distribution. SiC-particles impede the movement of dislocations glide and climb, resulting in an upper yield stress.

4. Conclusions The AlMg–SiCp composite is fabricated by vortex method with not secondary chemical reactions and particles distributed homogeneously. Vortex method has produced an excellent composite compared with a composite manufactured conventionally. Basically, molten metal is delivered via refractory feeder tip directly into the steel tube with SiCp preheated at 470 °C, where it solidifies and eliminate solidification defects.

Microstructural examination of composites produced under optimum condition mentioned above shows that distribution of particle reinforcement is homogeneous and products for secondary chemical reactions on the SiCp/matrix interface were not observed.

Al–Mg + SiC composite show an Al–Mg solid solution rich in aluminum and film aluminum oxide were not observed.

Thus a vortex technique is an excellent method to obtain composites with a homogeneous microstructure which allows us improves mechanical properties.

Acknowledgements The authors would like to thank Dr. O. Flores (Department of Materials, ICF-UNAM) for his technical support. Supported by DGAPA-UNAM (PAPIIT-IN105708).

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]

Fig. 3. Engineering stress–strain curves for Al–Mg unreinforced matrix alloy and Al–Mg/SiCp composite.

J. Hashim, J Mater Process Technol 123 (2002) 251–257.

P. Rohatgi, J Minerals 43 (1991) 10–15.

S.R. Nutt, R.W. Carpenter, Mater Sci Eng A 75 (1985) 169–174.

R. Mitra, V.S. Chalapathi Rao, R. Maiti, M. Chakraborty, Mater. Sci. Eng. A 379 (2004) 391–400.

C.G. Levi, G.J. Abbaschian, R. Mehrabian, Metall Trans. 9 (1978) 697–703.

N.E. Bekheet, Mater Design 23 (2002) 153–159.

S.V. Nair, J.K. Tien, R.C. Bates, Int. Met. Rev. 6 (1985) 275–290.

G.H. Paulino, J.H. Kim, J. Eng. Fract. Mech. 71 (2004) 1907–1912.

R. Warren, C.H. Anderson, Composite 15 (2) (1984) 101–107.

M. Gu, Z. Mei, Y. Jin, Z. Wu, Scripta Mater. 40 (9) (1999) 985–989.

M.J. Nathan, Mater. Sci. Lett. 8 (1999) 311–317.

H. Ribes, R. Da Silva, M. Suero, T. Bretheau, Mater. Sci. Technol. 6 (1990) 621–628.

G. Gonzalez, L. Salvo, M. Suero, G.L. Esperance, Scripta Metall. Mater.

33 (12) (1995) 1969–1973.

K.B. Lee, H. Kwon, Scripta Mater. 36 (8) (1997) 847–851.

H. Tadeo, I. Tsunemichi, Scripta Mater. 36 (2) (1997) 145–150.

A. Mehmet, G. Ferhat, Mater. Desig. 25 (2004) 209–217.

P. Villers, A. Prince, H. Okamoto, Handbook of Ternary Phase Diagrams, ASM International, Materials Park, OH, 1995, p. 2893.

Y. Brechet, Acta Metall. Mater. 39 (1991) 1781–1787.

J.R. Gomes, A.S. Miranda, D. Soares, A.E. Dias, L.A. Rocha, Ceram.

Trans. 114 (2000) 579–585.

T.S. Srivatsan, Mater. Trans. JIM 32 (1991) 473–479.

T.G. Durai, D.a.s. Karabi, D.a.s. Siddharta, Mater. Sci. Eng. A 445–446 (2007) 100–105.

P. Poddar, V.C. Srivastava, P.K. De, K.L. Sahoo, Mater. Sci. Eng. A 460–461 (2007) 357–364.