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

Microstructural characterization of extruded samples is discussed in terms of: (a) grain size; (b) Al₂O₃ nanoparticle reinforcement characteristics; and (c) crystallographic texture.

Microstructural analysis results revealed that there was no significant change in grain size for each nanocomposite compared to the corresponding monolithic alloy as shown in Table 1 and Figure 1. This suggested the inability of Al₂O₃ nanoparticles to serve as either nucleation sites or obstacles to grain growth during solid state cooling. Nearly equiaxed grains were observed in each monolithic material and derived nanocomposite.

X-ray diffraction (XRD) analysis revealed the presence of Al₁₂Mg₁₇ phase in the AZ31 based materials, and MgZn phase in the ZK60A based materials [21]. Al₂O₃ nanoparticle reinforcement as well as fine intermetallic distribution in the nanocomposites was reasonably uniform as shown in Figure 1. The reasonably uniform distribution of Al₂O₃ nanoparticles can be attributed to: (a) minimal gravity-associated segregation due to judicious selection of stirring parameters [22]; (b) good wetting of Al₂O₃ nanoparticles by the alloy matrix [23,24,25]; (c) argon gas disintegration of metallic stream [26]; and (d) dynamic deposition of composite slurry on substrate followed by hot extrusion. Selective area electron diffraction (SAED) in TEM mode revealed the partial reaction of Al₂O₃ nanoparticles with the magnesium alloy matrix to form MgO (where Al goes into solid solution and/or forms Mg-Al intermetallics) in each nanocomposite as shown in Figure 1 and listed in Table 2.

Texture results are listed in Table 3 and shown in Figure 2. In monolithic AZ31, the dominant textures in the transverse and longitudinal directions were (1 0 −1 0) and (1 0 −1 1) [and (0 0 0 2)], respectively. In the AZ31/Al₂O₃ nanocomposite, the dominant textures in the transverse and longitudinal directions were (1 0 −1 0) and (1 0 −1 1), respectively. Unlike monolithic AZ31, the AZ31/Al₂O₃ nanocomposite did not exhibit (0 0 0 2) dominant texture in the longitudinal direction. In monolithic ZK60A and derived nanocomposite, the dominant texture in the transverse and longitudinal directions was (1 0 −1 1).

The results of microhardness measurements are listed in Table 1. Each nanocomposite exhibited higher hardness than the corresponding monolithic alloy (by 30% and 15% for the AZ31/Al₂O₃ and ZK60A/Al₂O₃ nanocomposites, respectively). This was consistent with earlier observations made on Mg/Al₂O₃, AZ31/C₆₀ and AZ31/MWCNT nanocomposites [14,27,28]. The increase in hardness of the nanocomposites in the present study can be attributed to: (a) reasonably uniform distribution of harder Al₂O₃ nanoparticles in the matrix and (b) higher constraint to localized matrix deformation during indentation due to the presence of nanoparticles [14,15,27].

The overall results of ambient temperature tensile testing of the extruded materials are shown in Table 4 and Figure 3. The yield strength, ultimate strength, failure strain and work of fracture (WOF) of AZ31/Al₂O₃ nanocomposite were each significantly higher compared to monolithic alloy (by 19%, 21%, 113% and 162%, respectively). The WOF was determined by computing the area under the stress-strain curve. In the case of ZK60A/Al₂O₃ nanocomposite, the yield strength was decreased (by 4%) compared to monolithic alloy, while ultimate strength, failure strain and WOF were each significantly increased (by 13%, 170% and 200%, respectively).

Given the limited slip system activation in the HCP unit cell based structure of the AZ31 matrix at room temperature, the enhancement of tensile strength in the AZ31/Al₂O₃ nanocomposite can be first attributed to crystallographic texture difference between the nanocomposite matrix and monolithic alloy. In comparison of crystallographic texture, monolithic AZ31 exhibited (0 0 0 2) dominant texture in the longitudinal direction as listed in Table 3 and shown in Figure 2, unlike the AZ31/Al₂O₃ nanocomposite. For this (0 0 0 2) basal plane texture, basal slip is made most difficult due to the high critical resolved shear stress (CRSS) for slip based on the 0° angle between the (0 0 0 2) basal plane and the vertical axis as shown in Figure 2 [29,30]. However, non-basal slip was activated during deformation at room temperature due to basal plane alignment along the vertical (force) axis [31]. The tensile strength increase in the AZ31/Al₂O₃ nanocomposite compared to monolithic AZ31 can be next attributed to better 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 [14,15,32,33]; (b) Orowan strengthening mechanism [32,33]; and (c) load transfer from matrix to reinforcement [14,32]. The same supporting attributes apply to the increase (by 13%) of ultimate strength in the ZK60A/Al₂O₃ nanocomposite compared to monolithic alloy. In the case of ZK60A/Al₂O₃ nanocomposite, the yield strength was decreased (by 4%) compared to monolithic alloy. This can be attributed to relatively increased leaching of Zr from the ZK60A matrix by Al₂O₃ nanoparticles [34].

In each nanocomposite, the failure strain increase compared to monolithic alloy can be attributed to the presence and reasonably uniform distribution of ceramic nanoparticles [27,35]. It has been shown in previous studies that the 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 neighborhood of crack tip [27,35]. WOF quantified the ability of the material to absorb energy up to fracture under load [36]. In each nanocomposite, the significantly increased WOF compared to monolithic alloy show the potential for using the nanocomposite in damage tolerant design.

The overall results of ambient temperature compressive testing of the extruded materials are shown in Table 5 and Figure 4. The yield strength, ultimate strength and WOF of the AZ31/Al₂O₃ nanocomposite were each higher compared to monolithic alloy (by 5%, 5% and 11%, respectively), while failure strain of the AZ31/Al₂O₃ nanocomposite was not significantly decreased (by 4%) compared to monolithic alloy. In the case of the ZK60A/Al₂O₃ nanocomposite, the yield strength was decreased (by 10%) compared to monolithic alloy, while ultimate strength, failure strain and WOF were each increased (by 7%, 15% and 26%, respectively).

In each nanocomposite, the increase in strength can be attributed to similar better known factors (pertaining to reinforcement) as listed earlier regarding tensile strength increase. In the case of AZ31/Al₂O₃ nanocomposite, the increase in strength can be attributed to the nanocomposite not exhibiting (0 0 0 2) dominant texture in the longitudinal direction (unlike monolithic AZ31) as discussed earlier regarding tensile strength increase. In the case of the ZK60A/Al₂O₃ nanocomposite, the yield strength was decreased (by 10%) compared to monolithic alloy. This can again be attributed to the relatively increased leaching of Zr from the ZK60A matrix by Al₂O₃ nanoparticles [34].

In the case of the ZK60A/Al₂O₃ nanocomposite, failure strain increase compared to monolithic alloy can be attributed to the same reason as listed for tensile failure strain increase. In each nanocomposite, the increased WOF compared to monolithic alloy show the potential for using the nanocomposite in damage tolerant design.

In both the AZ31/Al₂O₃ nanocomposite and monolithic AZ31, 0.2% TYS was about double the 0.2% CYS (2.08 and 1.85, respectively, where yield strength anisotropy or YSA = 0.2% TYS/0.2% CYS). This can be attributed generally to {1 0 1 −2} < 1 0 1 −1 > -type twinning being more easily activated along the c-axis (of the HCP unit cell in Figure 2) during c-axis tension rather than c-axis compression [37,38]. With the c-axis of the HCP unit cell oriented 90° to the force axis (extrusion direction) in each case, {1 0 1 −2} < 1 0 1 −1 > -type twinning was less easily activated during tensile testing, when the c-axis was under compression.

In both the ZK60A/Al₂O₃ nanocomposite and monolithic ZK60A, 0.2% TYS was about 1.99 and 1.96 times (almost double) the 0.2% CYS, respectively. Here, the tensile/compressive YSA was approximately 2 despite the common crystallographic texture exhibited where {1 0 1 −2} < 1 0 1 −1 > -type twinning was activated along the c-axis of the hexagonal close packed (HCP) unit cell in Figure 2 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 [37,38]. Here, the tensile/compressive yield stress anisotropy can be attributed generally to half the strain rate used (less strain hardening) in compressive testing compared to tensile testing.

In this study, AZ31 (nominally 2.50–3.50 wt.% Al, 0.60–1.40 wt.% Zn, 0.15–0.40 wt.% Mn, 0.10 wt.% Si, 0.05 wt.% Cu, 0.01 wt.% Fe, 0.01 wt.% Ni, balance Mg) rod supplied by Asianovena Pte Ltd (Singapore) and ZK60A (nominally 4.80–6.20 wt.% Zn, 0.45 wt.% Zr, balance Mg) supplied by Tokyo Magnesium Co. Ltd. (Yokohama, Japan) were used as the matrix material. AZ31 rod and ZK60A block were sectioned to smaller pieces. All oxide and scale surfaces were removed using machining. All surfaces were washed with ethanol after machining. Al₂O₃ nanopowder (50 nm size) supplied by Baikowski (Japan) was used as the reinforcement phase.

Monolithic AZ31 and ZK60A were cast using the disintegrated melt deposition (DMD) method [21,39]. This involved heating AZ31 or ZK60A sections 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 melt 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 25l pm. The disintegrated melt was subsequently deposited onto a metallic substrate located 500 mm from the disintegration point. Ingots of 40 mm diameter were obtained following the deposition stage. To form the AZ31/Al₂O₃ and ZK60A/Al₂O₃ nanocomposites, Al₂O₃ nanoparticle powder was isolated by wrapping in Al foil of minimal weight (<0.50 wt.% with respect to AZ31 or ZK60A matrix weight) and arranged on top of the AZ31 or ZK60A blocks, with all other DMD parameters unchanged. All ingots were sectioned into billets. All billets were machined to 36 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.

Microstructural characterization studies were conducted on metallographically polished monolithic and nanocomposite extruded samples to determine grain characteristics. Hitachi S4300 Field-Emission SEM was used. Image analysis using Scion software was carried out to determine the grain characteristics. Thin foils were prepared from the nanocomposite extruded samples using disc punch-out and ion-milling for (regarding localized effects): (a) nanoparticle distribution as well as (b) nanoparticle-matrix reactivity observation (based on selected area electron diffraction (SAED)) using transmission electron microscopy (JEOL JEM 3010 TEM with 300KeV accelerating voltage). Regarding SAED, nanopowder samples of Al₂O₃ dispersed in ethanol were also prepared by droplet application onto holey carbon film mounted on copper grids followed by drying. All k-space measurements (k) from SAED patterns were manually obtained and converted to crystallographic lattice d-spacings (d) based on d = 2π/k. Goniometer 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 (regarding globalised effects). All d-spacings from SAED and goniometer XRD analysis were matched with corresponding d-spacings in the JCPDS database available in the Shimadzu LAB-X XRD-6000 diffractometer operating software to determine all phases present.

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.

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 a MTS 810 machine equipped with an axial extensometer with a crosshead speed set at 0.254 mm/min.

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 a MTS 810 machine with 0.005 min⁻¹ strain rate.

Monolithic AZ31 and the AZ31/Al₂O₃ nanocomposite can be successfully synthesized using the DMD technique followed by hot extrusion.

Compared to monolithic AZ31, tensile strength of the AZ31/Al₂O₃ nanocomposite was enhanced. Compared to monolithic AZ31, compressive strength of the AZ31/Al₂O₃ nanocomposite was slightly increased. This can be commonly attributed to: (a) AZ31/Al₂O₃ nanocomposite not exhibiting (0 0 0 2) dominant texture in the longitudinal direction (unlike monolithic AZ31); and (b) well known factors pertaining to reinforcement.

Compared to monolithic AZ31, tensile failure strain of the AZ31/Al₂O₃ nanocomposite was enhanced and the compressive failure strain was similar. This can be commonly attributed to reasonably uniform distribution of Al₂O₃ nanoparticles in the AZ31/Al₂O₃ nanocomposite.

Compared to monolithic AZ31, the AZ31/Al₂O₃ nanocomposite exhibited significantly high increment in tensile WOF and increment in compressive WOF.

Monolithic ZK60A and the ZK60A/Al₂O₃ nanocomposite can be successfully synthesized using the DMD technique followed by hot extrusion.

Compared to monolithic ZK60A, yield and ultimate strengths (in tension and compression) of the ZK60A/Al₂O₃ nanocomposite were decreased and increased, respectively. The decrease in yield strength can be attributed to the relatively increased leaching of Zr from the ZK60A matrix by Al₂O₃ nanoparticles. The increase in ultimate strength can be attributed to well known factors pertaining to reinforcement.

Compared to monolithic ZK60A, tensile and compressive failure strain of the ZK60A/Al₂O₃ nanocomposite was significantly enhanced and enhanced, respectively. This can be attributed to the presence and reasonably uniform distribution of Al₂O₃ nanoparticles.

Compared to monolithic ZK60A, the ZK60A/Al₂O₃ nanocomposite exhibited significantly high increment in tensile WOF and increment in compressive WOF.