Effects of in situ TiB2 particle on microstructures and mechanical properties of AZ91 alloy

Effects of in situ TiB2 particle on microstructures and mechanical properties of AZ91 alloy

Journal of Alloys and Compounds 494 (2010) 94–97 Contents lists available at ScienceDirect Journal of Alloys and Compounds journal homepage: www.els...

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Journal of Alloys and Compounds 494 (2010) 94–97

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom

Effects of in situ TiB2 particle on microstructures and mechanical properties of AZ91 alloy S.-L. Zhang, Y.-T. Zhao ∗ , G. Chen, X.-N. Cheng, C.-J. She, X.-Y. Wang, D.-N. Wu School of Materials Science and Engineering, Jiangsu University, Zhenjiang, Jiangsu 212013, PR China

a r t i c l e

i n f o

Article history: Received 24 August 2009 Received in revised form 13 October 2009 Accepted 14 October 2009 Available online 22 October 2009 Keywords: AZ91 alloy In situ TiB2 particle Microstructure Mechanical properties

a b s t r a c t The effects of in situ TiB2 particle fabricated from Al–K2 TiF6 –KBF4 system via Direct Melt Reaction (DMR) technology on microstructure and mechanical properties of AZ91 alloy are investigated. The results indicate that the sizes of ␣-Mg sosoloids are decreased and the morphologies are orbed with in situ TiB2 particle addition; the amounts of ␤-Mg17 Al12 are decreased and the morphologies are granulated. The mechanical properties tests show that the tensile strength  b , elongation ı and HB hardness are 221.70 MPa, 3.87% and 75.49 kgf/mm2 , which are 1.19, 0.98 and 1.22 times to that of without Al–TiB2 addition, respectively. The effect mechanisms are discussed. © 2009 Elsevier B.V. All rights reserved.

1. Introduction Magnesium alloys are greatly demanded in automobile industry because of their low density, high specific strength and stiffness, superb damping capacity, electromagnetic shielding capacities, excellent machinability and good castability [1–4]. Among the various magnesium alloys developed, AZ91 is the most common commercial magnesium alloy, which is widely used cast magnesium alloy not only in automobile industry but also in other structural components assembly due to its best combination of mechanical properties and castability [5–8]. However, the applications of this alloy in critical conditions such as power train components, engine blocks, etc. are restricted because of the coarsening ␤-Mg17 Al12 phase with body-centered-cubic (bcc) structure present in the AZ91 alloy. The body-centered cubic (bcc) structure of Mg17 Al12 is incompatible with the hexagonal-close-packed (hcp) structure of magnesium matrix, which leads to the fragility of the Mg–Mg17 Al12 interface. In addition, Mg17 Al12 phase is relatively soft and has poor strength [9]. These greatly influence the mechanical properties of AZ91 alloy. In order to further improve the strength and resistance to effects of heat, some researching work [10–15] has been done by micro-alloying technology, thermomechanical treatment, powder metallurgical methods, etc. Bayani and Saebnoori [16] have report that the addition of rare earth elements reduces Mg17 Al12

∗ Corresponding author. Tel.: +86 511 8791947; fax: +86 511 8791947. E-mail addresses: [email protected], [email protected] (Y.-T. Zhao). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.10.103

phase volume fraction and consequent decrease of the brittle Mg/Mg17 Al12 interface which was the main reason for weak thermal properties of the alloy at rather high temperatures. Ref. [17] shows that with 1.0 wt.% calcium and 0.5 wt.% strontium adding, the cast AZ91 magnesium alloy had fine grain sizes of 19 ␮m, the tensile strength and elongation at room temperature reaches 250 MPa and 3.5%, respectively and the strength at an elevated temperature of 448 K on developed AZ91 magnesium alloy is also drastically larger than that on precious cast magnesium alloys. In the present work, the effects of in situ TiB2 particle fabricated from Al–K2 TiF6 –KBF4 system via Direct Melt Reaction (DMR) technology on microstructure and mechanical properties of AZ91 alloy are investigated and the effect mechanisms are discussed. 2. Experimental procedures The experimental procedures were as following by two-step method. Firstly, the Al–TiB2 intermediate alloy was prepared via Direct Melt Reaction (DMR) technology from Al–K2 TiF6 –KBF4 system. The nominal compositions of K2 TiF6 and KBF4 powder (average size 20 ␮m) are listed in Table 1. When the temperature of the aluminum melt was 1123 K, the dehumidified K2 TiF6 and KBF4 powder (where the mole ratio is 1:2.2) with the weight ratio 20 wt.% to the total aluminum melt were added into the aluminum molten with a campanulate graphite plunge. After 20 min, the melt was degassed by hexachloroethane degasifying agent, deslagged, and then was poured into a copper mould. After cooled to the room temperature in air, the specimens were processed into nubs with linear cutting machine for the next procedure using. Secondly, the AZ91 alloy ingot was melt in a crucible under CO2 and SF6 gas protection atmosphere. When the molten temperature was 1003 K, the Al–TiB2 intermediate alloy (preheated to 523 K) was added into the AZ91 molten and the molten temperature was heated up to 1003 K. In addition, the mechanical stirring process was carried out for the Al–TiB2 intermediate alloy distributing in the molten. After 10 min, the molten was slagged off and refined. When the molten temperature was 993 K, the molten was poured into a copper mould. After cooled to the room

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Table 1 Nominal compositions of the AZ91 ingots, K2 TiF6 and KBF4 powder (wt.%). AZ91 K2 TiF5 KBf4 powder

Al 8.92 Fe2 O3 ≤0.05 Fe2 O3 ≤0.03

Zn 0.65 H2 O ≤0.15 H2 O ≤0.10

Mn 0.22 K2 TiF5 Balance KBF4 Balance

Fe 0.0001

Cu 0.0041

Si 0.0432

Ni 0.0008

Be 0.0012

Cl 0.0012

Mg Balance

Fig. 1. Sizes of the specimens for mechanical properties testing.

temperature in air, the specimens were cut two parts, the small one was used for XRD and SEM analysis and the large one for mechanical properties testing. The sizes of the specimens for mechanical properties testing were shown in Fig. 1. X-ray diffractometer (Dmax2500PC) using Cu K␣ radiation was used to the determine phase component of the as-prepared specimens; scanning electron microscopy (SEM, JOEL-7001F) was used to analysis the microstructures of the asprepared specimens and the tensile fracture surfaces. Tensile properties tests of the specimens were carried out at room temperature by a computer-controlled electronic tensile testing machine (DWD-200) at a strain rate 1.67 × 10−4 s−1 according to the ASTM E8 standard. The values of the tensile properties reported are the average of three tests at each condition, respectively.

3. Results and discussion

Fig. 2. XRD pattern of the Al–TiB2 Al–K2 TiF6 –KBF4 system.

intermediate alloy fabricated from

3.1. Preparations of Al–TiB2 intermediate alloy

Al3 Ti phases according to Refs. [18–20], respectively. However, the peak of Al3 Ti phase is very weak due to the excessive additions of KBF4 powder. This means some chemical reactions take place in the aluminum melt as following:

Fig. 2 shows the XRD pattern of the Al–TiB2 intermediate alloy fabricated from Al–K2 TiF6 –KBF4 system. It consists of Al, TiB2 and

3K2 TiF6 + 13Al = TiAl3 + 3KAlF4 + K3 AlF6

(1)

Fig. 3. SEM microstructures of the Al–TiB2 intermediate alloy fabricated from Al–K2 TiF6 –KBF4 system: (a) distributions of in situ TiB2 particles in aluminum matrix; (b) morphologies of in situ TiB2 particles; (c) EDS spectrum graph of the E area in (b).

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2KBF4 + 3Al = AlB2 + 2KAlF4

(2)

TiAl3 + AlB2 = TiB2 + 4Al

(3)

The total reactions can be expressed as 2KBF4 + 3K2 TiF6 + 12Al = 5KAlF4 + K3 AlF6 + TiB2

(4)

Fig. 3 shows the SEM microstructures and the energy dispersive X-ray spectroscopy (EDS) spectrum graph of the Al–TiB2 intermediate alloy fabricated from Al–K2 TiF6 –KBF4 system. As shown in Fig. 3a, the in situ TiB2 particles are distributed with disperse-state in the aluminum matrix. Fig. 3b shows the high magnified SEM pattern of the A area in Fig. 3a. It indicates that the morphologies of the TiB2 particles are in near-ball-shape or olive-shape; the size reaches micro-nanometer standard, and most is less than 0.2 ␮m. Fig. 3c shows the EDS spectrum graph. It reveals that the Al, Ti and B three elements. The size of the in situ TiB2 particle reaches nanometer, the analyses area is possibly included the aluminum matrix. This leads to the existence of aluminum element in the EDS spectrum graph. 3.2. Microstructures of the AZ91 alloy The cast structures of AZ91 alloy are consist of ␣-Mg and ␤Mg17 Al12 according to Mg–Al binary alloy phase. Fig. 4 shows the effects of the as-prepared in situ TiB2 phases on the microstructures of the AZ91 alloy. As shown in Fig. 4a, the ␣-Mg and ␤-Mg17 Al12 phases are coarse, and the ␣-Mg presented in irregular shape, the ␤-Mg17 Al12 phase are distributed in discontinuance meshy state along the boundary of the ␣-Mg phase. With the Al–TiB2 intermediate alloy addition, the ␣-Mg phases are circled and refined, the ␤-Mg17 Al12 phases are refined and mostly changed into massiveshape or particle-shape, as shown in Fig. 4b. These are due to the in

Table 2 The mechanical properties of the as-prepared AZ91 alloy fabricated with and without Al–TiB2 addition. Al–TiB2

Tensile strength (MPa)

Elongation (%)

Brinell hardness (kgf/mm2 )

Without addition With addition

186.25 221.70

3.94 3.87

62.13 75.49

situ TiB2 phase introduced and the effect mechanisms are discussed as following. TiB2 phase owns the hexagonal crystal structure with a = 0.3030 nm and c = 0.3230 nm, which is compatible with the hexagonal structure of magnesium matrix. The TiB2 phases are response of the heterogeneity nucleation substratum. Theses contribute to the amount increasing of the heterogeneity nucleation for ␣-Mg phase and the ␣-Mg phases are refined. The variations of the morphologies and distributions of the ␤-Mg17 Al12 phase are due to the refining of the ␣-Mg phases. When the ␣-Mg phases are refined, the amounts of the crystal boundaries are increased, and the volume fractions of the solutes are decreased. These lead to the amount reductions of the ␤-Mg17 Al12 phase per unit crystal boundary area and the uniform distributions after the solidifying ending. 3.3. Mechanical properties Table 2 gives the mechanical properties of the as-prepared AZ91 alloy fabricated with and without Al–TiB2 addition. It reveals that the mechanical properties of the as-prepared AZ91 alloy are increased than that of without Al–TiB2 addition. With Al–TiB2 addition, the tensile strength  b , elongation ı and HB hardness reach 221.70 MPa, 3.87% and 75.49 kgf/mm2 , which are 1.19, 0.98 and

Fig. 4. Effects of the as-prepared in situ TiB2 phases on the microstructures of the AZ91 alloy: (a) and (b) without Al–TiB2 additions, OM and SEM graphs; (c) and (d) with Al–TiB2 additions OM and SEM graphs.

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Fig. 5. SEM images of the tensile fracture surface of the as-prepared AZ91 alloy with and without Al–TiB2 additions: (a) without Al–TiB2 additions; (b) with Al–TiB2 additions.

1.22 times to that of without Al–TiB2 addition, respectively. The improvements of AZ91 mechanical properties are attributed from TiB2 phase which has the same the crystalline structure as the matrix, hexagonal crystal structure. Fig. 5 shows the SEM images of the tensile fracture surface of the as-prepared AZ91 alloy with and without Al–TiB2 additions, respectively. It can be clearly seen that the appearances of the fracture surfaces are different in size and depth with and without Al–TiB2 additions. As shown in Fig. 5a, the tensile fracture surface show the obvious cleavage fracture characteristics, the brittlement ␤Mg17 Al12 phases are clearly seen in the fracture surface and become the main crack fountainhead, and the observed dimples are shallower in depth and larger in size, and parts are linked along the boundaries. With Al–TiB2 addition, the brittlement ␤-Mg17 Al12 and ␣-Mg phases are refined, the sizes of the crystal grain are decreased. Thus, the ratio of the surface area and volume is increased, and the surface tension increased. These countercheck the expansion of the deforming area along the boundary of the crystal grain. As shown in Fig. 5b, the observed dimples are deeper in depth and small in size than that of without Al–TiB2 addition. The appearances of the tearing deformation inside the crystal are clearly observed, the cleavage fracture characteristics are weakened. The fracture mechanisms are the hybrid fracture mechanisms: the plastic fracture and cleavage fracture. 4. Conclusions The effects of in situ TiB2 particle fabricated from Al–K2 TiF6 –KBF4 system via Direct Melt Reaction (DMR) technology on microstructure and mechanical properties of AZ91 alloy are investigated by XRD, SEM. The results indicate that with in situ TiB2 particle addition, the sizes of ␣-Mg sosoloids are decreased, and the morphologies orbed; the amounts of ␤-Mg17 Al12 are decreased and the morphologies are granulated. The mechanical properties tests show that the values of the tensile strength  b , elongation ı and HB hardness are increased than that of the initial AZ91 alloy, and reached 221.70 MPa, 3.87% and 75.49 kgf/mm2 , respectively. The tensile fracture surface observations show that without that the appearances of the fracture surfaces are different in size and depth with and without Al–TiB2 additions. Without Al–TiB2 addition, the tensile fracture surfaces show the obvious cleavage fracture characteristics. With Al–TiB2 addition, the appearances of the tearing deformation inside the crystal are clearly observed, the cleavage fracture characteristics are

weakened, and the fracture mechanisms are the hybrid fracture mechanisms: the plastic fracture and cleavage fracture. Acknowledgements This work was supported by the Specialized Research Fund Project for the Doctoral Program of Higher Education of China (No. 20070299004), Jiangsu Province High-tech Research Project (No. BG2007030), Jiangsu Higher Education Institutions Natural Science Foundation Major Fundamental Research Program Project (07KJA43008), Jiangsu University Research Funding for Advanced Scholars (No. 1283000349) and Jiangsu University Undergraduate Practice-Innovation Training Project (No. 1201220022). References [1] H. Friedrich, S. Schumann, Journal of Materials Processing Technology 117 (3) (2001) 276–281. [2] A. Stalmann, W. Sebastian, H. Friedrich, S. Schumann, K. Dröder, Advanced Engineering Materials 3 (12) (2001) 969–974. [3] V.N. Chuvil’deev, T.G. Nieh, M.Yu. Gryaznov, V.I. Kopylov, A.N. Sysoev, Journal of Alloys and Compounds 378 (1–2) (2004) 253–257. [4] S. Lee, Y.-H. Chen, J.-Y. Wang, Journal of Materials Processing Technology 124 (1–2) (2002) 19–24. [5] K. Máthis, J. Gubicza, N.H. Nam, Journal of Alloys and Compounds 394 (1–2) (2005) 194–199. [6] P. Cavaliere, P.P. De, Marco, Journal of Materials Processing Technology 184 (1–3) (2007) 77–83. [7] Y. Guangyin, S. Yangshan, D. Wenjiang, Scripta Materialia 43 (11) (2000) 1009–1013. [8] S. Begum, D.L. Chen, S. Xu, A.A. Luo, International Journal of Fatigue 31 (4) (2009) 726–735. [9] B. Hossein, E. Saebnoori, Journal of Rare Earths 27 (2) (2009) 255–258. [10] J. Chen, Z. Chen, H. Yan, F. Zhang, K. Liao, Journal of Alloys and Compounds 461 (1–2) (2008) 209–215. [11] Z. Zhao, Q. Chen, Y. Wang, D. Shu, Materials Science and Engineering A 515 (2009) 152–161. [12] C. Scharf, A. Ditze, A. Shkurankov, E. Morales, C. Blawert, W. Dietze, K.U. Kainer, Advanced Engineering Materials 7 (12) (2006) 1134–1142. [13] C. Wen, Y. Yamada, K. Shimojima, M. Mabuchi, M. Nakamura, T. Asahina, T. Aizawa, K. Higashi, Materials Transactions JIM (Jpn Inst Met) 41 (9) (2000) 1192–1195. [14] Z. Zhang, D.L. Chen, Scripta Materialia 54 (7) (2006) 1321–1326. [15] A. Maltais, D. Dubé, M. Fiset, G. Laroche, S. Turgeon, Materials Characterization 52 (2) (2004) 103–119. [16] H. Bayani, E. Saebnoori, Journal of Rare Earths 27 (2) (2009) 255–258. [17] K. Hirai, H. Somekawa, Y. Takigawa, K. Higashi, Materials Science and Engineering A 403 (2005) 276–280. [18] Swanson, Tatge, Natl. Bur. Stand. (U.S.), Circ. 539,111 (1953). [19] J.T. Norton, H. Blumenthal, S.J. Sindeband, Transactions on AIME 185 (1949) 739. [20] J. Braun, M. Ellner, B. Predel, Zeitschrift fur Metallkunde 85 (1994) 855.