Adding TiC Nanoparticles to Magnesium Alloy ZK60A for Strength/Ductility Enhancement

Paramsothy, Muralidharan; Chan, Jimmy; Kwok, Richard; Gupta, Manoj

doi:https://doi.org/10.1155/2011/642980

Journal of Nanomaterials

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Research Article | Open Access

Volume 2011 | Article ID 642980 | https://doi.org/10.1155/2011/642980

Adding TiC Nanoparticles to Magnesium Alloy ZK60A for Strength/Ductility Enhancement

Muralidharan Paramsothy,¹Jimmy Chan,²Richard Kwok,²and Manoj Gupta¹

Academic Editor: Theodorian Borca-Tasciuc

Received10 Mar 2011

Accepted29 Jun 2011

Published05 Sept 2011

Abstract

ZK60A nanocomposite containing TiC nanoparticles was fabricated using solidification processing followed by hot extrusion. The ZK60A nanocomposite exhibited similar grain size to monolithic ZK60A and significantly reduced presence of intermetallic phase, reasonable TiC nanoparticle distribution, nondominant (0 0 0 2) texture in the longitudinal direction, and 16% lower hardness than monolithic ZK60A. Compared to monolithic ZK60A (in tension), the ZK60A nanocomposite simultaneously exhibited higher 0.2% TYS, UTS, failure strain, and work of fracture (WOF) (+13%, +15%, +76%, and +106%, resp.). Also, compared to monolithic ZK60A (in compression), the ZK60A nanocomposite exhibited lower 0.2% CYS (−17%) and higher UCS, failure strain, and WOF (+11%, +29%, and +34%, resp.). The beneficial effect of adding TiC nanoparticles on the enhanced tensile and compressive response of ZK60A is investigated in this paper.

1. Introduction

Magnesium alloys based on the Mg-Zn combination (such as ZK60A from the Mg-Zn-Zr system) are well known for their precipitation hardening characteristics during ageing [1]. Here, MgZn’ (or metastable β₁’ phase) forms as rods parallel to the c-axis of the HCP unit cell, while metastable β₂’ phase forms as discs parallel to the (0 0 0 2) basal plane of the HCP unit cell [2–6]. The β₁ phase has been reported to have a monoclinic structure similar to that of Mg₄Zn₇ in aged Mg-8 wt.% Zn alloy [2]. The yield strength of quasicrystalline particle reinforced Mg-Zn-Y and Mg-Zfn-Y-Zr magnesium alloys was observed to increase with the volume fraction of the quasicrystalline phase [7]. This was based on the strengthening effect of the quasicrystalline particles [7]. Due to low particle-matrix interfacial energy, icosahedral particles in the Mg-Zn-Y alloy have been observed to be stable against coarsening during elevated temperature deformation [7]. ZK60A is commonly used in structural applications based on the following characteristics: (a) high strength and ductility after T5 aging, (b) good creep resistance, (c) poor arc weldability due to hot-shortness cracking, and (d) excellent resistance weldability. Al₂O₃ nanoparticles were added to ZK60A recently using disintegrated melt deposition (DMD) [8, 9] followed by hot extrusion and heat treatment. In the study, the Al₂O₃ nanoparticles were agglomerated and not well dispersed in the ZK60A matrix [10]. This resulted in (1) tensile/compressive strength of ZK60A increasing (without significant loss in ductility) in the presence of 1.0 vol.% Al₂O₃ nanoparticles and (2) tensile/compressive ductility increasing (without significant loss in strength) in the presence of 1.5 vol.% Al₂O₃ nanoparticles [10]. It was also observed that Zr and Zn were leached out of the ZK60A matrix by the agglomerates of Al₂O₃ nanoparticles [10]. Carbon nanotubes (CNTs) were also added to ZK60A recently using DMD [8, 9] followed by hot extrusion. Here, CNTs were not agglomerated but reasonably well distributed in the ZK60A matrix [11]. This enabled (1) simultaneous increase in tensile strength and ductility of ZK60A and (2) significant increase in compressive ductility (with significant decrease in compressive strength) of ZK60A [11]. Also, the intermetallic phase(s) precipitation was reported to be possibly regulated at nanoscale in this nanocomposite [11]. Dissolved Zn segregation at the liquid-SiC nanoparticle interface in a cast ZK60A/SiC nanocomposite which enabled nanoscale MgZn₂ precipitation has also been discussed in detail [12].

Ti and C are not known to react actively with molten Mg. This is favourable concerning metal-matrix composite processing for structural applications where particle-matrix interfacial reactions leading to inferior mechanical properties are undesirable. Regarding TiC reinforced metal-matrix composites, molten Al-Mg alloys were infiltrated at 900°C into TiC preforms with flowing argon [13]. Wetting of TiC substrates by the molten Al-Mg alloys was investigated in this work [13]. It was observed that (1) selective Al₄C₃ formation was present at the matrix-preform interface and (2) TiAl₃ traces were present in the Al-based matrix [13]. Also, a preform containing elemental powders of Ti and C (where TiC was formed in situ) was initially formed [14]. Molten magnesium alloy AZ91D was pressurelessly infiltrated into this Ti-C preform and tensile properties of the composite were compared to monolithic AZ91D [14]. Here, the in situ formed TiC reinforcement enhanced the tensile strength especially at higher temperatures [14]. The strain hardening exponent (n) of the AZ91D/TiC composite was higher at 0.71–0.82 compared to that of monolithic AZ91D being 0.11–0.32 (for tensile deformation carried out at 423–723 K in each case) [14]. The mechanical properties of hot pressed W were improved based on the addition of La₂O₃ and TiC [15]. Here, the W matrix was strengthened due to the reinforcement particles pinning down the grain boundaries and inhibiting grain growth during sintering [15]. TiC particles exhibited good interfacial characteristics (for effective load transfer) with the adjacent W matrix while La₂O₃ particles were beneficial for sintering and densification of composites [15]. The collective strengthening effect of La₂O₃ and TiC particles on W was better than that of either La₂O₃ or TiC [15]. It was recently reported that TiC formed in situ when CNT was added to pure Ti powder and the mixture was consolidated using spark plasma sintering followed by hot extrusion [16]. Here, the titanium matrix nanocomposite exhibited significantly higher yield and ultimate strengths without considerably compromising ductility compared to monolithic Ti [16]. Open literature search has revealed that no successful attempt has been made to simultaneously increase tensile strength and ductility of ZK60A magnesium alloy with TiC or any other Ti-based nanoparticles, using a high volume production spray-deposition-based solidification processing technique.

Accordingly, one of the primary aims of this study was to simultaneously increase tensile strength and ductility of ZK60A magnesium alloy with TiC nanoparticles. Another aim of the present study was to evaluate the compressive properties of ZK60A/TiC magnesium alloy nanocomposite. Disintegrated melt deposition (DMD) [8, 9] followed by hot extrusion was used to synthesize the ZK60A/TiC magnesium alloy nanocomposite.

2. Experimental Procedures

2.1. Materials

In this study, ZK60A (nominally 4.80–6.20 wt.% Zn, 0.45 wt.% Zr, balance Mg) supplied by Tokyo Magnesium Co. Ltd. (Yokohama, Japan) was used as the matrix material. ZK60A block was sectioned to smaller pieces. All oxide and scale surfaces were removed using machining. All surfaces were washed with ethanol after machining. TiC nanoparticles (98+% purity, 30–50 nm size) supplied by Nanostructured and Amorphous Materials Inc. (Tex, USA) was used as the reinforcement phase.

2.2. Processing

Monolithic ZK60A was cast using the DMD method [8, 9]. This involved heating ZK60A blocks to 750°C in an inert Ar gas atmosphere in a graphite crucible using a resistance heating furnace. The crucible was equipped with an arrangement for bottom pouring. Upon reaching the superheat temperature, the molten slurry was stirred for 2.5 min at 460 rpm using a twin blade (pitch 45°) mild steel impeller to facilitate the uniform distribution of heat. The impeller was coated with Zirtex 25 (86% ZrO₂, 8.8% Y₂O₃, 3.6% SiO₂, 1.2% K₂O and Na₂O, and 0.3% trace inorganics) to avoid iron contamination of the molten metal. The melt was then released through a 10 mm diameter orifice at the base of the crucible. The melt was disintegrated by two jets of argon gas oriented normal to the melt stream and located 265 mm from the melt pouring point. The argon gas flow rate was maintained at 25 dm³ min⁻¹. The disintegrated melt slurry was subsequently deposited onto a metallic substrate located 500 mm from the disintegration point. An ingot of 40 mm diameter was obtained following the deposition stage. To form the ZK60A/1.5 vol.% TiC nanocomposite (see Figure 1), TiC nanoparticle powder was isolated by wrapping in Al foil of minimal weight (<0.50 wt.% with respect to ZK60A matrix weight) and arranged on top of the ZK60A blocks, with all other DMD parameters unchanged. All deposited ingots were sectioned into billets. All billets were machined to 35 mm diameter and hot extruded using 20.25 : 1 extrusion ratio on a 150 ton hydraulic press. The extrusion temperature was 350°C. The billets were held at 400°C for 60 min in a furnace prior to extrusion. Colloidal graphite was used as a lubricant. Rods of 8 mm were obtained.

2.3. Microstructural Characterization

Microstructural characterization studies were conducted on metallographically polished monolithic and nanocomposite extruded samples to determine grain characteristics as well as nanoparticle reinforcement distribution. Hitachi S4300 Field-Emission SEM was used. Image analysis using Scion software was carried out to determine the grain characteristics. XRD studies were conducted using CuK_α radiation (λ = 1.5406 Å) with a scan speed of 2°/min in an automated Shimadzu LAB-X XRD-6000 diffractometer to determine intermetallic phase(s) presence and dominant textures in the transverse and longitudinal (extrusion) directions.

2.4. Hardness

Microhardness measurements were made on polished monolithic and nanocomposite extruded samples. Vickers microhardness was measured with an automatic digital Shimadzu HMV Microhardness Tester using 25 gf-indenting load and 15 s dwell time.

2.5. Tensile Testing

Smooth bar tensile properties of the monolithic and nanocomposite extruded samples were determined based on ASTM E8M-05. Round tension test samples of 5 mm diameter and 25 mm gauge length were subjected to tension using an MTS 810 machine equipped with an axial extensometer with a crosshead speed set at 0.254 mm/min. Fractography was performed on the tensile fracture surfaces using Hitachi S4300 FESEM.

2.6. Compressive Testing

Compressive properties of the monolithic and nanocomposite extruded samples were determined based on ASTM E9-89a. Samples of 8 mm length (l) and 8 mm diameter (d) where l/d = 1 were subjected to compression using an MTS 810 machine with 0.005 min⁻¹ strain rate. Fractography was performed on the compressive fracture surfaces using Hitachi S4300 FESEM.

3. Results

3.1. Macrostructural Characteristics

No macropores or shrinkage cavities were observed in the cast monolithic and nanocomposite materials. No macrostructural defects were observed for extruded rods of monolithic and nanocomposite materials.

3.2. Microstructural Characteristics

Microstructural analysis results revealed that grain size was statistically unchanged in the case of nanocomposite as shown in Table 1 and Figure 2(a). Unlike monolithic ZK60A, intermetallic phase(s) were not detected in the nanocomposite by X-ray diffraction as shown in Figure 2(b). The distribution of individual TiC nanoparticles present in the nanocomposite (at the grain boundary and within the grain) was reasonably uniform as shown in Figures 2(c) and 2(d).

(a)

(b)

(c)

(d)

Texture results are listed in Table 2 and shown in Figure 3. In monolithic and nanocomposite materials, the dominant texture in the transverse and longitudinal directions was (1 0 −1 1).

3.3. Hardness

The results of microhardness measurements are listed in Table 1. The nanocomposite exhibited lower hardness than the monolithic material.

3.4. Tensile Behavior

The overall results of ambient temperature tensile testing of the extruded materials are shown in Table 3 and Figure 4(a). The strength, failure strain, and work of fracture (WOF) of ZK60A/1.5 vol.% TiC were significantly higher compared to monolithic ZK60A. The WOF was determined by computing the area under the stress-strain curve up to the point of fracture. The fractured surface of all extruded materials exhibited mixed (ductile + brittle) mode of fracture as shown in Figures 5(a) and 5(b).

(a)

(b)

(a)

(b)

(c)

(d)

3.5. Compressive Behavior

The overall results of ambient temperature compressive testing of the extruded materials are shown in Table 4 and Figure 4(b). The 0.2% CYS of ZK60A/1.5 vol.% TiC was lower compared to monolithic ZK60A. The UCS, failure strain, and work of fracture (WOF) of ZK60A/1.5 vol.% TiC were significantly higher compared to monolithic ZK60A. The fractured surface of ZK60A/1.5 vol.% TiC appeared smoother than that of monolithic ZK60A as shown in Figures 5(c) and 5(d).

4. Discussion

4.1. Synthesis of Monolithic ZK60A and ZK60A/TiC Nanocomposite

Synthesis of monolithic and nanocomposite materials, the final form being extruded rods, was successfully accomplished with (a) no detectable metal oxidation and (b) no detectable reaction between graphite crucible and melts. The inert atmosphere used during DMD was effective in preventing oxidation of the Mg melt. No stable carbides formed due to reaction with the graphite crucible.

4.2. Microstructural Characteristics

Microstructural characterization of extruded samples is discussed in terms of (a) grain and intermetallic phase(s) characteristics and (b) TiC nanoparticle reinforcement distribution.

Nearly equiaxed grains were observed in monolithic material and nanocomposite as shown in Table 1 and Figure 2(a). Grain size was statistically unchanged in the case of nanocomposite, suggesting the inability of TiC nanoparticles to serve as either nucleation sites or obstacles to grain growth during solid state cooling. Intermetallic phase(s) were not observed within the resolution of XRD (less than 2% by volume) as well as FESEM in the nanocomposite. This indicated that the intermetallic particle size in the nanocomposite was significantly lower compared to that of monolithic ZK60A. It is well known that crystalline nanoparticles can have insufficient Bragg diffraction planes needed for exhibiting strong diffraction peaks in goniometer-based XRD [17]. This resulted in the corresponding absence of strong diffraction peaks pertaining to intermetallic phase(s) in the XRD scan. The intermetallic phase(s) precipitation was possibly regulated at nanoscale due to the presence of (a) dissolved Al and/or (b) well-dispersed TiC nanoparticles. In the case of (a), dissolved Al possibly altered the intermetallic phase stabilities in the ZK60A matrix. This is similar to free Si from SiC nanoparticles being possibly responsible for altered intermetallic phase stabilities in Mg-Zn alloys as recently reported [12]. In the case of (b), dissolved Zn possibly segregated at the liquid-TiC nanoparticle interface enabling intermetallic phase manipulation at the nanoscale. This is similar to possible dissolved Zn segregation at the liquid-SiC nanoparticle interface enabling nanoscale MgZn₂ precipitation as recently reported [12]. With a reasonably uniform 1.5 vol.% TiC distribution throughout the ZK60A matrix, the nanoparticle-matrix interface area was ample for effectively regulated segregation of 4.80–6.20 wt.% Zn (or 1.21–1.59 vol.% Zn) as nanoscale Mg-Zn precipitates. This was similar to that reported recently for ZK60A/1.0 vol.% CNT nanocomposite [11].

The reasonably uniform distribution of TiC nanoparticles can be attributed to (a) minimal gravity-associated segregation due to judicious selection of stirring parameters [8], (b) good wetting of TiC nanoparticles by the alloy matrix [18–20], (c) argon gas disintegration of metallic stream [21], and (d) dynamic deposition of composite slurry on substrate followed by hot extrusion. Similar reasonably uniform distribution of CNT nanoparticles in magnesium alloy ZK60A has also been recently reported [11]. In the nanocomposite, no reaction products based on: (a) Mg (or Al) and TiC (such as Mg₂C₃, Al₄C₃, or Al-Ti-based intermetallic in this case) and (b) Al and Zr (such as fine Al-Zr-based intermetallic in this case) having more than 2% by volume were detected using X-ray diffraction analysis.

4.3. Mechanical Behavior

4.3.1. Hardness

A significant decrease in microhardness (−16%) was observed in the nanocomposite compared to monolithic material as listed in Table 1. There was lower constraint to localized matrix deformation in the nanocomposite compared to monolithic material. This was despite the reasonably uniform distribution of harder TiC nanoparticles in the nanocomposite (see Figures 2(c) and 2(d)). The decrease in hardness can be primarily attributed to significantly reduced precipitate size of intermetallic phase(s) in the matrix of the nanocomposite compared to monolithic material (see Figure 2(b)).

4.3.2. Tensile and Compressive Behavior

Strength
The tensile and compressive strengths of monolithic material and nanocomposite are listed in Tables 3 and 4 (and shown in Figures 4(a) and 4(b)), respectively. 0.2% TYS and UTS were enhanced (+13% and +15%, resp.) in ZK60A/1.5 vol.% TiC compared to monolithic material. The tensile strength increase in ZK60A/1.5 vol.% TiC compared to monolithic ZK60A was despite the significantly reduced precipitate size of intermetallic phase(s) in the matrix of the nanocomposite compared to monolithic material (see Figure 2(b)). This increase in tensile strength can be attributed to the overall positive effect derived from well-known factors pertaining to reinforcement such as (a) dislocation generation due to elastic modulus mismatch and coefficient of thermal expansion mismatch between the matrix and reinforcement [22–25], (b) Orowan strengthening mechanism [24–26], and (c) load transfer from matrix to reinforcement [22, 24].
Regarding compressive strengths, 0.2% CYS and UCS of ZK60A/1.5 vol.% TiC were lower and higher (−17% and +11%), respectively, compared to monolithic ZK60A. The compressive stress detected at any given strain was lower for ZK60A/1.5 vol.% TiC compared to monolithic ZK60A as shown in Figure 4(b). This was despite the factors pertaining to reinforcement (as just described in the paragraph before this). This lower compressive strength in ZK60A/1.5 vol.% TiC compared to monolithic ZK60A can be attributed to the overall negative effect derived from (a) significantly reduced precipitate size of intermetallic phase(s) in the matrix of the nanocomposite compared to monolithic material and (b) compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic as illustrated in Figure 6. The dominant presence of (a) Mg-Zn-based nanoscale rod intermetallic (lying parallel to the c-axis of the HCP unit cell) and (b) Mg-Zn based nanoscale disc intermetallic (lying parallel to the basal plane of the HCP unit cell) in Mg-Zn systems is well known [1, 27]. The compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic induced a significantly lower limit on the strengthening factors pertaining to reinforcement (as just described in the paragraph before this). This was dissimilar to that reported recently for ZK60A/1.0 vol.% CNT nanocomposite, where CNT buckling in the ZK60A matrix was the reason compressive stress detected at any given strain was lower for ZK60A/1.0 vol.% CNT compared to monolithic ZK60A [11].
In monolithic ZK60A, 0.2% TYS was about 1.27 times the 0.2% CYS. Here, near tensile/compressive yield stress isotropy was present despite half the strain rate used (less strain hardening) in compressive testing compared to tensile testing. This can be attributed to {1 0 1 −2}<1 0 1 −1>-type twinning being activated along the c-axis of the HCP unit cell in Figure 3 (see Table 2 also) with comparatively similar ease in both tension and compression along the c-axis, based on the 45° angle between the c-axis and the vertical axis [20, 28, 29]. In the case of ZK60A/1.5 vol.% TiC, 0.2% TYS was about 1.74 times the 0.2% CYS despite the similarity in crystallographic texture compared to monolithic ZK60A. Here, the significant tensile/compressive yield stress anisotropy can be attributed to compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic in ZK60A/1.5 vol.% TiC as illustrated in Figure 6. The Mg-Zn-based nanoscale rod/disc intermetallic is prone to buckling followed by fracture within the ZK60A matrix during compressive deformation unlike during tensile deformation. This was similar to that reported recently for ZK60A/1.0 vol.% CNT nanocomposite [11].

Failure Strain
The tensile and compressive failure strains of monolithic material and nanocomposite are listed in Tables 3 and 4 (and shown in Figures 4(a) and 4(b)), respectively. Compared to monolithic material, tensile and compressive failure strains were enhanced by 76% and 29%, respectively in ZK60A/1.5 vol.% TiC. The failure strain increase in ZK60A/1.5 vol.% TiC compared to monolithic ZK60A can be attributed to the following factors pertaining to reinforcement: (a) presence and reasonably uniform distribution of TiC nanoparticles [30, 31], (b) reduction in size of intermetallic phase(s) [32], and (c) compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic as illustrated in Figure 6 (regarding compressive failure strain only). In the case of reasonably uniform distribution of TiC nanoparticles, it has been shown in previous studies that nanoparticles provide sites where cleavage cracks are opened ahead of the advancing crack front. This (1) dissipates the stress concentration which would otherwise exist at the crack front and (2) alters the local effective stress state from plane strain to plane stress in the neighbourhood of the crack tip [30, 31]. In the case of reduction in size of intermetallic phase(s), redistribution of smaller intermetallic phase(s) (compare between predominantly aggregated type and dispersed type) can assist in improving ductility [32]. Regarding compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic, such nanoscale buckling within the ZK60A matrix aids in dispersing localized stored energy during compressive deformation. This allows ZK60A/1.5 vol.% TiC to globally absorb relatively large amounts of strain energy during compressive deformation [11, 33]. Here, Mg-Zn-based nanoscale rod/disc intermetallic buckling within the ZK60A matrix is a compressive toughening mechanism. This was dissimilar to that reported recently for ZK60A/1.0 vol.% CNT nanocomposite, where CNT buckling in the ZK60A matrix was one of the reasons compressive failure strain was higher for ZK60A/1.0 vol.% CNT compared to monolithic ZK60A [11].
Tensile fracture behaviour of both monolithic material and nanocomposite was mixed (ductile + brittle) as shown in Figures 5(a) and 5(b). However, the tensile fractured surface of the nanocomposite had (a) higher occurrence of smaller dimple-like features and (b) absence of microcracks compared to that of monolithic material. The tensile cavitation resistance was lower and the microcrack formation resistance was higher in the nanocomposite compared to monolithic material. For ZK60A/1.5 vol.% TiC, compressive fracture behavior based on viscoplastic flow [34] was relatively more ductile (smoother fracture surface exhibited) compared to monolithic ZK60A as shown in Figures 5(c) and 5(d). This relatively more ductile compressive fracture behavior can be attributed to increase in shear band spacing [34, 35].

Work of Fracture
The tensile and compressive work of fracture (WOF) of monolithic material and nanocomposite are listed in Tables 3 and 4 (and illustrated in Figures 4(a) and 4(b)), respectively. WOF quantified the ability of the material to absorb energy up to fracture under load [36]. Compared to monolithic material, tensile and compressive WOF were enhanced by 106% and 34% (resp.) in ZK60A/1.5 vol.% TiC. The significantly high increments in tensile and compressive WOF exhibited by ZK60A/1.5 vol.% TiC show its potential to be used in damage tolerant design.

5. Conclusions

(1)Monolithic ZK60A and ZK60A/1.5 vol.% TiC nanocomposite can be successfully synthesized using the DMD technique followed by hot extrusion.(2)Compared to monolithic ZK60A, tensile strength of ZK60A/1.5 vol.% TiC was enhanced. This can be attributed to the overall positive effect derived from well-known factors pertaining to reinforcement. Compared to monolithic ZK60A, compressive strength of ZK60A/1.5 vol.% TiC was generally decreased. The 0.2% CYS and UCS of ZK60A/1.5 vol.% TiC were lower and higher than that of monolithic ZK60A, respectively. This can be attributed to the overall negative effect derived from (a) significantly reduced precipitate size of intermetallic phase(s) in the matrix of the nanocomposite compared to monolithic material and (b) compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic in the nanocomposite.(3)Compared to monolithic ZK60A, tensile and compressive failure strains of ZK60A/1.5 vol.% TiC were significantly enhanced. This can be attributed to the following factors pertaining to reinforcement: (a) presence and reasonably uniform distribution of TiC nanoparticles, (b) reduction in size of intermetallic phase(s), and (c) compressive shear buckling of Mg-Zn-based nanoscale rod/disc intermetallic (regarding compressive failure strain only).(4)Compared to monolithic ZK60A, ZK60A/1.5 vol.% TiC exhibited significantly high increments in tensile and compressive WOF.

Acknowledgments

Authors wish to acknowledge National University of Singapore (NUS) and Temasek Defence Systems Institute (TDSI) for funding this research (TDSI/09-011/1A and WBS no. R265000349).

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Copyright © 2011 Muralidharan Paramsothy et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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