Si₃N₄ Nanoparticle Addition to Concentrated Magnesium Alloy AZ81: Enhanced Tensile Ductility and Compressive Strength

Abstract

This study is aimed at understanding the tensile ductility and compressive strength-enhancing dual function of nanoparticles in a concentrated magnesium alloy (AZ81) nanocomposite. Si₃N₄ nanoparticles were selected for reinforcement purposes due to the known affinity between magnesium and nitrogen. AZ81 magnesium alloy was reinforced with Si₃N₄ nanoparticles using solidification processing followed by hot extrusion. The nanocomposite exhibited similar grain size and hardness to the monolithic alloy, reasonable nanoparticle distribution, and nondominant (0 0 0 2) texture in the longitudinal direction. Compared to the monolithic alloy in tension, the nanocomposite exhibited higher failure strain (+23%) without significant compromise in strength, and higher energy absorbed until fracture (EA) (+27%). Compared to the monolithic alloy in compression, the nanocomposite exhibited similar failure strain (+3%) with significant increase in strength (up to +20%) and higher EA (+24%). The beneficial effects of Si₃N₄ nanoparticle addition on tensile ductility and compressive strength dual enhancement of AZ81 alloy are discussed in this paper.

1. Introduction

Silicon nitride nanoparticles in the shape of near-spheres and wires have been synthesized by chemical vapor deposition (CVD) [1, 2]. In the case of nanowires, the catalytic properties of metal nanoparticles were utilized during CVD to promote the directed silicon nitride growth at nanoscale [2]. On one hand, nanoparticles are active towards cells (or bioactive), where the specific biological function at cellular level can be disrupted, modified (negatively or positively), or promoted by the uncoated nanoparticles present [3]. On the other hand, silicon nitride nanoparticles have been coated and chemically stabilized prior to effective dispersion in a rubber matrix [4]. In the context of mechanical (crystallographic structure related) or functional (electronic structure related) properties, the function of nanoparticles in a metallic matrix is related to (a) nanoparticle-matrix reactivity and (b) nanoparticle distribution in the matrix. Magnesium alloys are an easily available lightweight and energy saving metallic matrix. In particular, the AZ (Aluminium-Zinc) series of magnesium alloys are characterized by (a) low cost, (b) ease of handling, (c) good strength and ductility, and (d) resistance to atmospheric corrosion [5]. These qualities enable the common use of AZ series magnesium alloys [5]. Regarding magnesium nanocomposites, the friction stir processing technique has been used to add SiO₂ nanoparticles to AZ61 [6]. Here, the tensile elongation at 350°C of selected composites was lower than 100% at 1×10⁻³ s⁻¹ strain rate but reached 350% and 420% at 1×10⁻² s⁻¹ and 1×10⁻¹ s⁻¹ strain rates, respectively. High strain rate superplasticity was exhibited taking into consideration the sufficiently uniform dispersion of SiO₂ nanoparticles. In another case, the superplasticity of Mg-Zn-Zr composite systems containing SiC microparticles or submicroparticles has been studied and attributed to (a) fine grain size (lesser twinning effects) and (b) crystallographic textural effects [7, 8]. Concerning solidification processed magnesium alloy nanocomposites, existing representative research literature indicates that (a) good nanoparticle distribution can be achieved in the magnesium matrix and (b) better mechanical properties (specifically ductility) can be achieved due to the addition of nanoparticles [9–13].

The physical attributes of second phase precipitates influence the mechanical properties of the alloy. Regarding concentrated AZ series magnesium alloys (e.g., AZ80 and AZ91), the Mg-Al second phase presence significantly contributes to strengthening. Solutionizing heat treatment (where the Mg-Al second phase goes into solution): (a) drastically removes the strengthening effect and (b) significantly contributes to ductility enhancement. Not much is known about the interplay between nanoscale reinforcement presence and second phase precipitation (which is especially relevant in concentrated magnesium alloy nanocomposites) concerning the effects on mechanical properties. Accordingly, this study is aimed at understanding the tensile ductility and compressive strength-enhancing dual function of nanoparticles in an AZ81 alloy nanocomposite. AZ81 was reinforced with Si₃N₄ nanoparticles using solidification processing (disintegrated melt deposition (DMD) [14, 15]) followed by hot extrusion.

2. Results and Discussion

2.1. Synthesis of Monolithic AZ81 and Derived AZ81/Si₃N₄ 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 the graphite crucible and melt. The inert atmosphere used during 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. 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.

2.2. Microstructural Characteristics

Microstructural analysis results revealed that grain size and aspect ratio remained statistically unchanged as shown in Table 1 and Figure 1(a). Distribution of nanoparticles and fine intermetallic particles in the nanocomposite was reasonably uniform as shown in Figure 1(b). Goniometer XRD analysis revealed the presence of Al₃Mg₂ and Al₁₂Mg₁₇ intermetallics in monolithic AZ81 and the derived nanocomposite. 3 crystallographic planes corresponding to Mg-Al second phase were detected in AZ81/Si₃N₄ as indicated by TEM SAED phase analysis results listed in Table 2 and shown in Figure 1(c). Texture results are listed in Table 3 and shown in Figure 2. In monolithic AZ81 and the derived nanocomposite, the dominant texture in the transverse and longitudinal directions was (1 0 1 1).

Table 1 

Results of grain characteristics and microhardness of monolithic AZ81 and derived nanocomposite. (Prior to tensile/compressive testing).

Table 2 

Phase detection results of AZ81/Si₃N₄ nanocomposite and Si₃N₄ nanoparticle reinforcement based on selected area electron diffraction (TEM). (Prior to tensile/compressive testing).

Table 3 

Texture results of monolithic AZ81 and derived nanocomposite based on X-ray diffraction (goniometer). (Prior to tensile/compressive testing).

(a)  monolithic alloy/nanocomposite

(b) nanocomposite

(c) nanocomposite

(d) Si3N4 nanoparticles

Figure 1 

(a) Representative SEM micrograph showing grain size in monolithic AZ81 and AZ81/Si3N4 nanocomposite. (b) Representative TEM micrograph showing the presence of individual nanoparticles (fine intermetallics inclusive) in AZ81/Si3N4 nanocomposite. Representative SAED diffraction patterns of (c) AZ81/Si3N4 nanocomposite and (d) Si3N4 nanoparticles. Linked with Table 3. (Prior to tensile/compressive testing).

Figure 2 

Schematic diagram showing texture of monolithic AZ81 and derived nanocomposite based on goniometer X-ray diffraction. In each case, vertical axis (dotted line) is parallel to extrusion direction. Each cell is made up of 2 HCP units having 1 common (0 0 0 2) basal plane. (Prior to tensile/compressive testing).

2.3. Tensile/Compressive Behavior

The overall results of ambient temperature tensile testing of the extruded materials are shown in Table 4 and Figure 3. The overall results of ambient temperature compressive testing of the extruded materials are shown in Table 5 and Figure 4. The tensile failure strain increase in the nanocomposite compared to the corresponding monolithic alloy can be attributed to the following factors (pertaining to reinforcement) (a) presence and reasonably uniform distribution of ceramic nanoparticles [16, 17] and (b) Si₃N₄ nanoparticle induced regulated precipitation of Mg-Al second phase as illustrated in Figure 5. Regarding factor (a), 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 neighbourhood of the crack tip [16, 17]. Regarding factor (b), dissolved Al possibly segregated at the liquid-Si₃N₄ nanoparticle interface enabling Mg-Al second 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 [13]. The Si₃N₄ nanoparticle was originally amorphous in the as-supplied form but adopted a crystalline structure in the nanocomposite as shown in Figures 1(d) and 1(c) (resp.) and listed in Table 2. With a reasonably uniform 1.5 vol% Si₃N₄ distribution throughout the AZ81 matrix, the nanoparticle-matrix interface area was insufficient for effectively regulated segregation of 7.80–9.20 wt% Al (or 5.15–6.13 vol% Al) as nanoscale Mg-Al precipitates. The Si₃N₄ nanoparticle regulated precipitation of Mg-Al second phase is illustrated in Figure 5(c). Generally, while the Si₃N₄ nanoparticle is active in nanoscale manipulation of the Mg-Al second phase, the second phase nanoparticles newly formed move away, allowing more second phase nanoparticles to be formed adjacent to the active Si₃N₄ nanoparticle. This occurs until the Si₃N₄ nanoparticle reacts with the Mg based matrix to form a larger Mg₃N₂ particle (Mg₃N₂ particle formation not shown in Figure 5(c)). 1 crystallographic plane corresponding to Mg₃N₂ was detected in AZ81/Si₃N₄ as indicated by TEM SAED phase analysis results listed in Table 2 and shown in Figure 1(c). Some of the second phase nanoparticles that moved away from the active Si₃N₄ nanoparticle initially to make room also begin to grow larger. When the rate of second phase nanoparticle formation from the active Si₃N₄ nanoparticle exceeds the rate of earlier-formed second phase nanoparticle growth, the fraction of second phase particles of relatively smaller size increases, as illustrated in Figure 5(b) compared to Figure 5(a). In this case of reduction in size of Mg-Al second phase, redistribution of smaller second phase (compare between predominantly aggregated type and dispersed type) assists in improving ductility [18].

Table 4 

Results of tensile testing of monolithic AZ81 and derived nanocomposite. Linked with Figure 3.

Table 5 

Results of compressive testing of monolithic AZ81 and derived nanocomposite. Linked with Figure 4.

Figure 3 

Representative tensile stress-strain curves of monolithic AZ81 and AZ81/Si3N4 nanocomposite.

Figure 4 

Representative compressive stress-strain curves of monolithic AZ81 and AZ81/Si3N4 nanocomposite.

Figure 5 

Schematic diagram showing second phase precipitation in (a) monolithic AZ81 and (b) AZ81/Si3N4 nanocomposite. (c) Si3N4 nanoparticle regulated second phase precipitation in AZ81/Si3N4 nanocomposite. The circular shape of the second phase precipitates is only for illustration purposes.

Any hypothetical strength increase in the nanocomposite (effect 1) can be attributed to the following well known factors (pertaining to reinforcement): (a) dislocation generation due to elastic modulus mismatch and coefficient of thermal expansion mismatch between the matrix and reinforcement [19–22], (b) Orowan strengthening mechanism [21–23], and (c) load transfer from matrix to reinforcement [19, 21]. However, any hypothetical strength decrease in the nanocomposite (effect 2) can be attributed to reduction in size of Mg-Al second phase (since there is no statistical change in grain size between AZ81 and AZ81/Si₃N₄). In this study, effects 1 and 2 equalled each other in tension, causing the AZ81/Si₃N₄ and AZ81 tensile stress-strain curves to overlap each other as shown in Figure 3. However, effect 1 significantly outweighed effect 2 in compression, causing the AZ81/Si₃N₄ stress-strain curve to be above the AZ81 stress-strain curve (maximum 65 MPa stress difference) as shown in Figure 4. This implied that the AZ81-Si₃N₄ nanoscale interface was relatively less adequate in tension (crack opening deformation) but relatively more adequate in compression (crack closing deformation), as is the case with a less adequate particle-matrix interface in a composite that is unable to fully accommodate the stress concentrating effects associated with crack opening (during tensile deformation) but otherwise functions satisfactorily in the absence of crack opening (during compressive deformation). In both AZ81/Si₃N₄ and monolithic AZ81, 0.2% TYS was about 2.04 and 2.44 times (>2.00) the 0.2% CYS, respectively. Here, the tensile/compressive yield stress anisotropy was despite the crystallographic texture exhibited where {101−2}⟨1011⟩-type twinning was activated along the 𝑐-axis of the 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 [24, 25]. 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. The tensile/compressive yield stress anisotropy was desirably lower in AZ81/Si₃N₄ compared to monolithic AZ81. This can be attributed to the AZ81-Si₃N₄ nanoscale interface being relatively weaker in tension but relatively stronger in compression.

3. Experimental Section

3.1. Materials

The materials used in this study are listed in Table 6. Essentially, AA1050 was added to AZ61 to increase the nominal aluminium content of AZ61 by 2 wt%. All alloy pieces were sectioned to smaller pieces. All oxide and scale surfaces were removed using machining. All surfaces were washed with ethanol after machining. Si₃N₄ nanoparticles were used as the reinforcement phase in AZ81 magnesium alloy.

Table 6 

Materials selected for this study.

3.2. Processing

The monolithic alloys were cast using the DMD method [14, 15]. This involved heating AZ61 and AA1050 pieces 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 located 265 mm from the melt pouring point. The argon gas flow rate was maintained at 25 lpm. 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 AZ81/1.5 vol% Si₃N₄ alloy nanocomposite, Si₃N₄ nanoparticle powder was isolated by wrapping in Al foil of minimal weight (<0.50 wt% with respect to the total matrix weight) and arranged on top of the alloy blocks with all other DMD parameters unchanged. 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.

3.3. Heat Treatment

Heat treatment was carried out on all extruded sections of AZ81 based material at 200°C for 1 hour using a resistance heating furnace. This selection of temperature and time was made in order to relax the AZ81 alloy matrix without recrystallization softening. The recrystallization temperature of AZ80 magnesium alloy (as the nearest matching commercial alloy in terms of composition following 10% cold work after 1 hour) is 345°C [5]. Prior to heat treatment, the sections were coated with colloidal graphite and wrapped in aluminum foil to minimize reaction with oxygen present in the furnace atmosphere.

3.4. Microstructural Characterization

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 300 KeV accelerating voltage). Regarding SAED, nanopowder samples of Si₃N₄ (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=1.5406Å) 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.

3.5. 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.

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

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

4. Conclusions

(1)Monolithic AZ81 and the AZ81/Si₃N₄ nanocomposite can be successfully synthesized using the DMD technique followed by hot extrusion.(2)In tension, the AZ81/Si₃N₄ nanocomposite stress-strain curve was overlapping the AZ81 curve. In compression, the AZ81/Si₃N₄ nanocomposite stress-strain curve was above the AZ81 curve. This can be attributed to the AZ81-Si₃N₄ nanoscale interface being relatively less adequate in tension (crack opening deformation) but relatively more adequate in compression (crack closing deformation). In both cases, reduction in size of Mg-Al second phase supported hypothetical decrease in strength.(3)The tensile failure strain increase in the AZ81/Si₃N₄ nanocomposite compared to monolithic AZ81 can be attributed to the following factors (pertaining to reinforcement): (a) presence and reasonably uniform distribution of ceramic nanoparticles and (b) Si₃N₄ nanoparticle induced regulated precipitation of Mg-Al second phase.(4)The AZ81/Si₃N₄ nanocomposite exhibited higher tensile and compressive energy absorbed until fracture (EA) compared to monolithic AZ81.

Acknowledgments

The 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 # R265000349).