Strength vs. Ductility: A Comparative Microstructural Analysis of Stir-Cast Commercially Pure Aluminum Alloy AA1100 Reinforced with Mg and Al₂O₃

The experimental materials employed in this research are a high-purity aluminum matrix bonded with Mg powder and Al₂O₃ nanoparticles to assess the distinct effects of these additions on microstructural development and mechanical integrity. Aluminum 1100 was selected as the base matrix due to its exceptional corrosion resistance and inherent ductility. However, its unalloyed tensile strength is typically low, measuring approximately 49 MPa. A high-purity aluminum wire (99.8% purity) was used as the matrix material to minimize interference from alloying elements during melting. It is a non-heat-treatable type of wrought alloy that is strengthened either by solid solution hardening or by strain hardening. Commercially pure Mg powder (99.9% purity) was added to the first sample to convert the AA1100 matrix to a binary Al-Mg system, which replicates the properties of the 5xxx series alloys. A 5% Mg additive will cause solid-solution strengthening, as Mg atoms substitute for aluminum in the lattice, creating internal strain that limits dislocation motion. The addition of 5% Mg to the first sample will cause the development of the intermetallic compound. In the case of the second sample, 5 wt.% of Al₂O₃ nanoparticles were added to make an MMNC. Ceramic Reinforcement: α-Al₂O₃ Al₂O₃ nanoparticles were purchased from Sigma-Aldrich (St. Louis, MO, USA), and the average initial particle size was 40 nm, and the purity was 99%. This particular form of Al₂O₃ is a ceramic strengthening material with high compressive strength, high hardness, and thermodynamic stability, which hinder dislocation motion and enable the material to undergo plastic deformation. Energy-Dispersive X-ray Spectroscopy (EDS) was used to determine the elemental composition of the starting materials and the developed samples. Some trace impurities, such as heavy elements like Iron (Fe) and silicon (Si, were present in low amounts. The chemical analysis of the two samples in this investigation is shown in Table 1.

Table 1. The experimental materials' chemical analysis

A pure graphite crucible and an electronic temperature controller were installed in an electric resistance furnace to achieve excellent thermal control. The furnace maintains a constant processing temperature of 720 ℃–800 ℃ to support the effective integration of reinforcements. A protective gas was used to prevent unnecessary oxidation and hydrogen absorption, which would otherwise pose risks when melting pure aluminum with reactive Mg. To avoid contamination of the 1100 alloy melt, a ceramic-coated steel stirrer was employed. Stirring was performed continuously at about 750 ℃ to obtain a homogeneous distribution of the reinforcement phases. Sufficient shear energy was provided to uniformly distribute the globular intermetallic particles and nanoparticles throughout the matrix. Kneading must continue until just before pouring into the permanent metallic die (usually mild steel), so that the reinforcement is not deposited due to density differences.

4.1 Optical microscopy analyses

As-received commercially pure AA1100 aluminum alloy wire is seen in the optical micrograph (OM) in Figure 1(a). One interesting aspect of this image is that it represents the feedstock raw material before melting and stir-casting were performed. Hence, it is included here only as background to the initial constituents for characterization, rather than as a direct structural control in the cast composites themselves. AA1100 is defined as commercially pure aluminum with a minimum content of 99.00%. The rest of that 1.0%, however, is metallurgically essential and is what we are, in fact, looking at in this micrograph: the α-Al solid solution. Due to its commercial purity, the matrix is very ductile and fairly soft. The white/bright specks, which are highly reflective and scattered throughout the matrix, are secondary intermetallic particles in the form of agglomerates [18]. Fe and Si are the main contaminants of AA1100. Since Fe has an almost insoluble nature in solid form at zero degrees (and it is insoluble at almost any degree), it is precipitated during solidification.

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Figure 1. The morphological features of: (a) As-received characterization of AA1100, (b) cast AA1100 with an addition of 5 wt.% of Mg powder, and (c) cast AA1100 with an addition of 5 wt.% of Al₂O₃ nanoparticles, obtained by optical microscopy

These are bright phases that are mainly insoluble Al-Fe and Al-Fe-Si intermetallic (alpha Al₁₂Fe₃Si or Al₃Fe). These harder, more apparent constituent particles commonly polish up in relief and are highly representative of the darker, removed aluminum. Incidence of cold working: the presence of particles is not the only salient feature of this morphology; the spatial distribution and the matrix's general texture are also salient. The received material has evidently undergone extensive plastic deformation, likely by cold rolling. The clearly visible horizontal striations or bands run across the entire field of sight. These are flow lines that reflect a direction of rolling. The initial equiaxed (spherical/polygonal) grains of the cast aluminum are severely crushed and have been extended along the mechanical work performed on them. The horizontal bands are not randomly filled with bright intermetallic particles; the latter follow these bands. The brittle intermetallic particles produced in the initial casting were broken apart and stretched along the rolling direction during rolling. The authors term such arranged groupings of particles "stringers". Although dislocations cannot be observed with an optical microscope (this requires Transmission electron microscopy (TEM)), the banding is grossly distorted, indicating a high degree of strain hardening. The driven feedstock wire exhibits a very high degree of deformation due to cold working. Importantly, however, this special strain-hardened microstructure with embedded twisted dislocations completely disappeared once the wire was liquid-melted at 720–800 ℃ for the liquid matrix of the composites. The actual grain boundaries of the α-Al matrix are not clearly resolved in this micrograph; instead, the etchant attacks regions under high stress and interfaces between constituent particles rather than the grain boundaries themselves [19]. To observe the grain structure of severely deformed aluminum, it is common to perform anodic oxidation (using Barker reagent) and then view it under polarized light.

An intriguing and complex microstructure is shown in Figure 1(b). When 5 wt.% of Mg is introduced to the commercially pure AA1100 base, the thermodynamic environment and physical appearance of the alloy change radically. In contrast to the highly directional, cold-worked matrix of wrought AA1100, the micrograph shows a highly heterogeneous structure with numerous precipitates. The continuous light background represents the α-Al matrix. Mg is highly soluble in aluminum at high temperatures (reaching about 14.9 wt.% at 450 ℃). Nonetheless, this does not increase beyond about 2 wt.% at room temperature; below this temperature, the solubility decreases dramatically. The system is supersaturated at room temperature, as the addition is 5% Mg. A substantial amount of that Mg is driven out of solid solution by the cooling rate or subsequent thermal processing (such as sintering or artificial aging), and the hardness and chemical reactivity of the matrix are significantly altered relative to the base (AA1100). Precipitation of the beta-Phase (Al₃Mg₂). The dark, highly concentrated networks and globular structures scattered across the matrix are the most noticeable features in this micrograph. The precipitation of the equilibrium beta-phase, which is usually known as Al₃Mg₂ (or sometimes denoted Mg₅Al₈), is driven by a strong thermodynamic driving force due to the supersaturation of Mg [20]. Nucleation is observed to occur both intergranular (along what would be grain or powder particle boundaries) and intragranular. The clustering indicates microsegregation of Mg in localized regions, which is very common when the alloy is cast at a medium cooling rate or processed by powder metallurgy, where diffusion rates dominate ultimate homogenization. The electronegativity of the beta-phase is very high when compared to the α-Al matrix. This is because the electrochemical difference renders the precipitates highly vulnerable to localized attack. Whilst AA1100 is commercially pure, the subsequent Fe and Si play an essential role when Mg is added. Si is highly affined with Mg. The small bags of Si found in the AA1100 base, too, will react selectively with the added Mg to create dark and hard intermetallic particles made up of Mg₂Si. Probably, these Mg₂Si constituents are some of the very discrete, darker, localized locations in the micrograph, and they tend to form in advance of the Al₃Mg₂ phase and furnish exceptional pinning locations to control the growth of the grains. Pitting and Preparation Artifacts. The most important thing when reading the particular micrograph is recognizing metallographic artifacts. The dark features appear haloed or as separate craters under a microscope. Because of their brittle nature, the intermetallic phases (Al₃Mg₂ and Mg₂Si) frequently crack and extrude from the softer α-Al matrix during grinding and mechanical polishing, leaving behind microscopic pits. Since the Mg-rich phases are quite anodic, standard etchants (such as Keller or Graff-Sargent) will etch these areas severely. The observed deep, dark globular masses suggest eroded-out pits in which the Mg-rich phases have been dissolved or worn deeper by the etchant, and not the literal physical removal of the precipitate [21].

It is important to recognize that the extreme variation of appearance in Figure 1(b) is partly a metallographic artifact: the deep, dark globular regions in the figure are pits that contain intermetallic inclusions which took up the etchant or were eroded out of the bulk material by it, and the mechanical polishing, and are formed because the intermetallic phase is very anodic and brittle; this intentional precipitation of secondary phases gives the alloy outstanding macroscopic ductility and strength, through providing the highly effective localized pinning sites which inhibit grain growth and govern the continuous solid-solution strain hardening.

Figure 1(c) depicts an informative micrograph of a metal matrix composite (MMC). The use of ceramic nanoparticles in a soft, commercially pure alloy such as AA1100 poses special thermodynamic and kinetic challenges, particularly regarding the Grain Structure and Matrix (α-Al). The main α-Al matrix forms the lighter, continuous background. The grain boundaries of such are not as smudged together as in wrought AA1100, resulting in rather equiaxed polygonal grains. The morphology is strongly indicative of the material having been worked along a path to allow it to form grains without having undergone any rigid subsequent directional deformation- probably a casting process (such as stir casting) or a route in a powder metallurgy (PM) process followed by sintering. The Agglomeration challenge is the most significant point identified in the black, one giant constituent area. 5% Al₂O₃ nanoparticles were added, but they cannot be resolved with an optical microscope (individual particles below 100 nm cannot be resolved). The observed dark regions are on the micron or tens-of-microns scale. As such, the particle agglomeration captured in this micrograph is extreme. Nanoparticles have a very large surface-area-to-volume ratio; hence, their surface energy is very high. Strong Van der Waals forces cause the Al₂O₃ particles to cluster naturally, thereby minimizing this thermodynamic instability. Moreover, Al₂O₃ is highly wetted with molten aluminum. In case such a composite was made by a liquid-state processing route (such as casting), the molten aluminum will tend to reject the ceramic particles rather than coat them, and squeeze them into large, separated agglomerates to reduce the liquid-solid interfacial area. They are not non-selectively scattered in the intragranular areas; rather, they are seriously segregated at the boundaries of the grains as well as on the interdendritic grounds. As solidification proceeds, the solid-liquid interface of growing α-Al grains physically drags the suspension of nanoparticle clusters along behind the freezing front [22]. The clusters are finally caught in the final areas to harden: the grain boundaries. If this were produced through powder metallurgy, this morphology suggests that the Al₂O₃ nanoparticles covered the edges of the much larger AA1100 matrix powders in the mixing stage, while being confined at the edges of the previous particle boundaries in the compaction and sintering stages. Two interacting phenomena occur here at 5%. At 5% addition, there is just enough volume fraction of the reinforcement to agglomerate, so that, at that point, the clusters start joining to form continuous networks along the grain boundaries. The hardness gradient between the very hard Al₂O₃ clusters and the very soft AA1100 matrix is enormous during the preparation of MMCs for optical microscopy. When subjected to grinding and polishing, these brittle agglomerates are highly likely to be ripped off by abrasive forces within the soft matrix, resulting in deep micro-voids and craters. The appearance of the arrowed area, which is rather dark and somewhat shadowy, and the distorted halo around the other large dark areas, is a sure indication that it beholds localized porosity and micro-voids resulting from particle pull-out, rather than merely the planar surface of the ceramic material itself.

4.2 Scanning Electron Microscopy morphology analyses

In Figure 2(a), the SEM micrograph is a high-fidelity image of the microstructure and localized phase distribution in an AA1100 alloy with 5% Mg. The main α-Al matrix forms the broad, dark gray background observed in Figure 2(a) and is relatively smooth. At room temperature, the equilibrium solubility of Mg in Al is extremely low (approximately 1.4 wt.%). The sample has greatly exceeded this solubility limit due to the addition of 5% Mg. This type of strengthening is important for strengthening the solid solution during straining of the aluminum crystal lattice, but the excess Mg must be tipped out of the matrix. These areas are relatively featureless in the matrix itself, as they should be in secondary electron (SE) imaging, unless heavy chemical etching has revealed a specific crystallographic plane. Secondary Phase Precipitation (Bright Regions): The most evident aspect of this micrograph is the highly interconnected, bright network that defines the grain boundaries as well as the interdendritic gallons of the microstructure. These networks are highly consistent with the beta-phase intermetallic, which is Al₃Mg₂ (or sometimes expressed as Al₈Mg₅). Since the 5% Mg addition is supersaturated, the surplus Mg is moved into the high-energy regions of the microstructure, the grain boundaries, as the microstructure cools or sinters [23]. This network can act as a physical hindrance to dislocation movement, significantly increasing the hardness and yield strength of the composite. Nevertheless, since Al₃Mg₂ is highly anodic relative to the α-Al matrix, a continuous grain-boundary network such as this leaves the alloy highly vulnerable to intergranular corrosion and stress corrosion cracking (SCC).

In the matrix shown in Figure 2(b), the striations are straight and parallel. They are not natural microstructural characteristics (such as slip bands); they are metallurgical scratches from polishing. During polishing of a composite containing extremely soft aluminum mixed with much harder Al₂O₃, the matrix becomes smeared, and any Al₂O₃ particles that come off the composite squeeze over the soft α-Al, causing these scratches. The particles have a very slight shadowing effect. As polishing proceeds, the soft Al₂O₃ erosion rate exceeds that of hard Al₂O₃, causing the ceramic particles to remain slightly above the surface. The Al₂O₃-reinforcement stage is the dark, blocky parts.

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Figure 2. The Scanning Electron Microscopy (SEM) morphology of AA100 aluminium alloy with additions: (a) with 5% Mg powder addition, and (b) with 5% Al₂O₃ nanoparticle addition

The particles are very jagged, with sharp edges. Sharp edges are considered a serious stress concentrator in an MMC. At such sharp corners in the matrix, under tensile or cyclic loading, cracks will sprout at most. Dark, deep cracks are observed within some of the larger Al₂O₃ particles. Al₂O₃ is extremely brittle. Such internal cracks must have been generated either at the mechanical compaction/mixing stage of synthesis (in the case of powder metallurgy) or by thermal stress on cooling (because of the huge difference in CTE between Al and Al₂O₃ [24]. The interface between nickel oxide particles and the AA1100 matrix governs the mechanical transfer of load. Wettability Issues: dark voids and gaps are visible on the peripheries of some particles. Aluminium and Al₂O₃ are thermodynamically antagonistic in wetting. Molten aluminium (had this been cast) or solid-state aluminium (had this been sintered) does not readily wet the ceramic surface without the addition of wetting agents (such as Mg). Interfacial Porosity vs. Pull-out: The black porosity at the particle edges may represent real interfacial porosity, as these regions are difficult to wet. They may also be of the pull-out type, in which pieces of the delicate Al₂O₃ peel off during grinding and polishing, leaving a hole.

4.3 Transmission electron microscopy morphology analyses

A deep examination of the AA1100 + 5% Mg system, using TEM micrographs in Figure 3, indicates a highly dynamic thermochemical environment during processing. Typical evidence of the great affinity of Mg to oxygen is the presence of nanoscale MgO in an Al-5% Mg alloy. Mg participates in the elimination of oxygen as an active oxygen scavenger. Oxygen can never be free during the synthesis of your alloy (through casting metallurgy). It may be dissolved in the melt, entrained by the processing atmosphere, or adsorbed as surface oxygen on raw materials. Since the Gibbs free energy of formation of MgO is rather negative under processing temperatures, the solute Mg will initially react with the available oxygen, forming small volumes of MgO crystals at the nanoscale in the aluminum matrix [25]. This is a formed-in-situ secondary phase that has far-reaching consequences for a material's structural integrity. Figure 3(a) is a general view of the physical interactions between the soft matrix and the freshly developed hard phases in a general manner. The giant, wave-like black sweeps on the left are contours of extinction (extinction contours). Under such a heavily alloyed system as that of Al-5% Mg, the difference in the atomic radius between Al and the large concentration of solute Mg puts heavy strain on the α-Al lattice. Moreover, in situ precipitation forms hard particles, and this (such as those observed in satin MgO clusters around) gives rise to sharp local stress fields. The TEM foil is naturally thin, so it bends into these stress concentrators, bending the crystal planes into diffraction alignment and producing these dramatic, dark bands.

This is evidence of a matrix experiencing significant internal stress directly linked to solid-solution and precipitation hardening. Dark and clumped structures are suspended in the lighter, electron-transparent matrix phases (the MgO and, possibly, the early-stage beta-phase Al₃Mg₂ intermetallic). Note that they are actually clumped into spherical-like accumulations rather than perfectly dispersed. Since they form by the reaction of trace oxygen with Mg, they tend to cluster around microscopic density features (such as grain boundaries) and voids, thereby lowering their surface energy. They are highly effective dislocation blockers because of their size (around 50–100 nm in the case of the clusters).

In Figure 3(b), a high-resolution image, the microstructural advantage of in situ formed precipitates over ex situ additions is evident. The picture shows a clear boundary between the α-Al₂O₃ matrix and an in situ MgO crystal with its (111) lattice planes. Since the MgO is growing internally as a result of the constituent atoms of the alloy itself, the interface is atomically clean. It contains no gaps, absorbed gas layers, or amorphous contamination layers that generally plague artificially mixed lattices. The edges of the aluminum that encounter MgO (111) planes are discontinuous. Such closeness will ensure an ideal metallurgical bond. When the load is applied to this alloy, the stress will flow smoothly from the ductile aluminum to the ultra-hard MgO crystal without the interface rupturing or separating. On a mechanical level, this HRTEM image indicates that the sample has hard, perfectly bonded, nanosized obstacles fully embedded in the matrix. Such dislocations, finding themselves in motion, will not be able to cut through this MgO crystal but must bend around it (Orowan bowing), imposing enormous demands on the amount of stress applied to them. This mechanism is quite consistent with the observed better strength profiles at the various fixed additions of 5 wt.% Mg [26].

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Figure 3. Transmission electron microscopy (TEM) images of AA1100 with 5% Mg addition at (a) 400 nm and (b) 2 nm scale bars

Once a material deforms through its line defects, known as dislocations, they move around. Moving up the crystal lattice of a pure soft metal, such as AA1100, encounters very little resistance from dislocations, which is why pure aluminum is highly malleable. An in-situ formed MgO nanoparticle is an impenetrable barrier. The particles are too hard and have an impeccable metallurgical bond with the matrix, so moving dislocations cannot cut or shear them. They are rather made to prostrate themselves around them. The fact that the dislocation bows changes into a back stress that diminishes considerably more with the applied force required to sustain deformation. The dislocation eventually circumnavigates the particle completely, leaving behind an Orowan ring and proceeding. This extra contribution is most likely the force driving the increase in yield strength. The Orowan-Ashby equation has been modified as the most rigorous calculation of the strength increment (in MPa), which is as follows [27]:

$\Delta \sigma_{\text {orowan }}=\frac{0.13 G b}{\lambda} \ln \left(\frac{r}{b}\right)$            (1)

where, G is the shear modulus of the aluminum matrix (26,000 MPa). b is the size of the Burgers vector of aluminum (approximately 0.286 nm). r is the mean radius of MgO nanoparticles (which you can measure directly from the TEM images; it is probably in the 2–10 nm range). λ is the interparticle separation, or the effective separation between the particles. The volume fraction (f) of MgO particles in the matrix and their radius (r) have a significant effect on the interparticle spacing. In the case of spherical particles, it is usually approximated as [28]:

$\lambda \approx r\left(\sqrt{\frac{2 \pi}{3 f}}\right)$            (2)

Oxygen traces are always present, even in samples melted under controlled conditions. It is present in the air in this neighborhood, either lodged in the microscopic pores of the crucible or in a very thin natural oxide coating on the added 5 wt.% Mg powder. Mg is an immensely stronger thermodynamic contender for oxygen than aluminum. When that 5% Mg powder was stirred into the high-temperature molten AA1100, the atoms of that trace oxygen were immediately "scavenged" by part of the molecules of the molten Mg. So, this process took place within the liquid metal:

Mg₂O₂ + O₂ = 2 MgO

As the sample cooled and solidified, the newly formed nanosized MgO molecules crystallized directly from the melt, becoming permanently fixed within the solid aluminum matrix.

In Figure 4(a), TEM is the most dynamic evidence contributing to the reasons why this composite is stronger than the base AA1100 alloy. What is important here is not just the particle and the matrix of peculiar features, but the bright, web-like network in the aluminum. The dark feature marked Al₂O₃ is much larger than 100 nm. On this scale, it is a sub-micron or micron-sized particle, either because the original powder was larger than nominal or because the agglomeration is unfortunately so dense during the mixing/melting stage. The key to this picture is the bright, sharp lines that make a cellular network in the image of the α-Al matrix. These high-density dislocation networks, in particular, create subgrain boundaries. There is a significant difference in the Coefficients of Thermal Expansion (CTE) between aluminum and Al₂O₃. When the composite cooled (or reached room temperature) after melting (or sintering), the aluminum matrix attempted to shrink much faster than the ceramic particles did. This created huge localized interface shear stresses. The aluminum lattice should have become harder to sustain this strain, and in fact, a little swarm of geometrically necessary dislocations (GNDs) was produced. These dislocations knotted and wove themselves into the vivid sub-grain frontiers that you observe here, virtually held in place by the particle of the Al₂O₃. This local work hardening around the ceramic particles correlates with an increase in the material's macroscopic yield strength [29].

The location where the aluminum matrix joins the ceramic addition is indicated by the exact point of contact, as shown in the High-Resolution TEM (HRTEM) image in Figure 4(b). This reveals the nature of (externally added) composite interfaces. The α-Al matrix in the upper right exhibits wonderfully resolved, parallel lattice fringes, indicating high crystallinity. But the diagonal band, marked by the label Al₂O₃, does not exhibit this long-range periodic order; it appears heavily speckled or amorphous. Such a phenomenon in aluminum MMCs is usually due to three causative factors.

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Figure 4. Transmission electron microscopy (TEM) images of AA1100 with 5% Al₂O₃ addition at (a) 200 nm and (b) 5 nm scale bars

To begin with, nanoparticles typically have a disordered, amorphous outer coating due to their high surface energy. Second, the grains may be an intermediate phase (such as in the possible transitional form of the spinel-like structure, like that of the so-called γ-Al₂O₃) as opposed to the perfectly crystalline corundum (in the form of α-Al₂O₃). Lastly, this can also be an artifact of TEM sample preparation, in which the hard ceramic surface is ruthlessly amorphized by aggressive ion milling, and the interface is physically sharp, with no epitaxy (in which atomic planes of the matrix smoothly slide around the particle). When molten aluminum is unable to wet pure Al₂O₃, the interface is strongly held together by mechanical interlocking, and the compressive residual stresses (due to the CTE mismatch above) lock the matrix around the particle [30]. Here, there is no acute interfacial voiding to observe, as much as possible; however, at the same time, the bond is chiefly physical, not a profound reaction zone of the chemical/metallurgical kind.

During the solidification of the composite (from the manufacturing temperature, say, the solidification temperature of α-Al, 660 ℃, to room temperature, 25 ℃), the two materials would become unevenly shrunken. The CTE of aluminum (α-Al) is about 23. 6 × 10^-6 K^-1; the CTE of Al₂O₃ (α-Al₂O₃) is 7.2 × 10^-6 K^-1. The Al₂O₃ particle will resist the inflexible Al₂O₃ particle. They can accommodate large volume mismatch by deforming the aluminum lattice, thereby forming a swarm of GNDs in the surrounding aluminum matrix. These distorted dislocations constitute the bright cell networks observed in the TEM micrographs. To obtain the resulting increase in strength, perform this in two operations.

Step A: Determine the enhanced dislocation density. The density of dislocations that are caused by this thermal mismatch is given by the following equation developed by Arsenault and Shi [31]:

$\Delta \rho=\frac{12 f \Delta \alpha \Delta T}{b d(1-f)}$             (3)

where,

•f = volume fraction of the particles (0.05).

•Δα = the difference between the CTE of the matrix and particle (16.4 × 10^-6 K^-1).

•ΔT = the temperature drop during cooling (660°C − 25°C = 635 ℃).

•b = the Burgers vector of the Al₂O₃ matrix (0.286 × 10^-9 m). (The original manuscript draft contained a typo here, listing the unit as 10^-6 K^-1. This has been corrected to meters to align with the accurate values you provided later in the text).

•d = the average diameter of the Al₂O₃ particles (in mm).

•12 = a geometric constant that takes the shape of a particle into account.

Step B: Determine the increment of yield strength(s) (Δσ CTED). On finding the number of new dislocations generated, using the Taylor relation to determine the amount they prevent plasticity (increasing yield strength) [32]:

$\Delta \sigma_{C T E}=\beta . G . b . \sqrt{\Delta \rho}$            (4)

where, β is the material constant 1.25 (FCC metals, such as aluminum, usually). G is the aluminum-matrix shear modulus (26,000 MPa). b is the Burgers vector 0.286 × 10^-9 m.

4.4 Energy-Dispersive X-ray Spectroscopy results analysis

The two reinforced aluminum alloys developed during the current research were also critically characterized using EDS spectra (Figure 5), which served as an in-depth tool for validating the sample synthesis. The base MatrixAA1100 Identification in both spectra shows a predominant aluminum peak at about 1.48 keV, associated with the primary matrix of the high-purity (99.8) AA1100, which was taken to melt. Both samples show small peaks at Si and Fe at 1.74 and 6.40 keV, respectively. These are inherent constituent intermetallic phases, such as Al3Fe or α-AlFeSi, as is the norm for the 1100 series specification. The AA1100 + 5 wt.% Mg Powder spectrum of Sample 1 in Figure 6(a) is confirmation of the success of converting the matrix to a binary Al-Mg system to replicate the characteristics of 5xxx series alloys. The Mg peak (1.25 keV) corresponds to a 5.09% concentration of Mg in the chemical analysis. The EDS results confirm that Mg atoms have replaced Al atoms in the lattice, causing internal strain that prevents dislocation movement and increases hardness. In addition to the matrix, EDS mapping proved very useful for detecting the beta-phase (Mg₂Al₃) intermetallic. These manifest as dark, bead-like chains along the grain boundaries, characteristic of a sensitized microstructure.

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Figure 5. Energy-Dispersive X-ray Spectroscopy (EDS) results: (a) AA1100 with 5% Mg, and (b) AA1100 with 5% Al₂O₃ nanoparticles

Table 4. Comparative summary of Energy-Dispersive X-ray Spectroscopy (EDS) data

Sample 2 in Figure 6(b) indicates the material to be an Al₂O₃ reinforced by ceramic particles. Oxygen (O) at a peak (0.52 keV) is the first sign of the reinforcement of the material with Al₂O₃. Compared with Sample 1, this alloy contains trace amounts of Mg. Evidence of ceramic phases: EDS lattice images of this sample showed that a portion of the solidifying material consisted of the ceramic phase, i.e., highly thermodynamically stable, hard nanoparticles of Al₂O₃ that were driven to the edges by growing Al crystals. This caused the particles to become trapped at interdendritic or grain boundaries rather than being perfectly trapped in the lattice [33]. Table 4 compares the EDS data for both samples.

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Figure 6. (a) X-ray diffraction (XRD) pattern of the AA1100 + 5 wt.% Mg powder, (b) XRD pattern of the AA1100 + 5 wt.% Al₂O₃ nanocomposite

4.5 X-ray diffraction analysis

The high-intensity peaks in the XRD plot in Figure 6(a) correspond to the α-Al matrix, which is predominantly face-centered cubic (FCC). Nevertheless, it is no longer pure AA1100; it is a highly alloyed matrix. XRD is used to determine the phases, including the primary α-Al matrix. Due to the very large atomic radius of Mg (1.60 A) compared to the atomic radius of aluminum (1.43 A), the Mg atoms will replace the aluminum atoms in the lattice. The replacement of an atom induces internal, local compressive strain fields within the matrix. The addition of 5% Mg powder can effectively transform the aluminum. The most important data from diffraction, including the standard 2Ɵ diffraction angles and the assigned (hkl) Miller indices of the primary matrix and all identified secondary phases (Al₃Mg₂, MgO, Mg₂Si, and α-Al₂O₃, Cu-K α radiation), are summarized in Table 5.

This enormous distortion of the lattice is usually observed in an XRD pattern in the form of broadening of the peaks and a slight decrease in the primary α-Al peaks in the final 2 theta position relative to a pure, unstrained aluminum baseline. This confirms the principle of solid-solution strengthening. Although this causes the primary peaks to shift, the intermetallic beta-phase, formed by the 5% Mg addition, is the crucial phase identified as Al₃Mg₂ (which, in some cases, is referred to as Mg₅Al₈). XRD is specifically used to confirm the presence of elemental transitions between Al and Mg, such as Al₃Mg₂. They are smaller secondary peaks between the primary α-Al peaks, whose positions depend on the volume fraction.

Table 5. 2Ɵ peaks and assigned (hkl) Miller indices for identified phases

Besides the main Al-Mg phases, there is evidence that highly reactive trace elements remain in their original crystallographic phases, which may manifest as small peaks or background differences in the XRD data. The commercially pure AA1100 base contains small amounts of impurities, namely Si and Fe. Si reacts selectively with Mg to form hard intermetallic particles because of Mg's high affinity for oxygen, and it acts as an active oxygen scavenger during meltdown. The resulting trace oxygen is then reacted with molten Mg to form nanoscale MgO crystals within the aluminum matrix. The resulting nanoscale crystals of the secondary phases extend, imposing far-reaching implications on the material's structural integrity. Figure 6(b) is a textbook sample of an MMNC diffractogram. The most pronounced peaks in the diffractogram in Figure 6(b) are directly linked to the FCC α-Al matrix. Unlike Mg, the ceramic Al₂O₃ nanoparticles do not dissolve into the aluminum crystal lattice, leaving the internal dimensions of the unit cell unchanged. Consequently, the primary α-Al diffraction peaks remain sharp and occur at their unaltered diffraction angles. The smaller, more intense secondary peaks between the dominant aluminum ones are the crystallographic signature of ceramic reinforcement. Al₂O₃ is characterized by high hardness, compressive strength, and thermodynamic stability. Its apparent crystallographic peaks indicate that the nanoparticles were not destroyed or dissolved by such high-temperature processing as 720 ℃–800 ℃ without losing their chemical character or breaking down into the melt. Then, during solidification, the aluminum crystals began to form and physically pushed these forcibly stable particles to the edges, entrapping them at the ends of interdendritic regions or at the borders of the grains, not perfectly within the lattice as they should have been, but at the ends of the interdendritic regions.

Metallurgically, the most important element of this XRD pattern is the relatively explicit lack of unaccounted peaks. Molten aluminum can no longer wet pure Al₂O₃; the interface is forced together by mechanical interlocking. This bond is primarily physical rather than a chemical or metallurgical reaction zone. Since the interface is fixed by a purely mechanical grip rather than by chemical interactions, the XRD pattern reveals a clear absence of new intermetallic reactions, consistent with a brittle character. The trend is neatly divided into two phases: α-Al and Al₂O₃. This two-period XRD pattern would confirm the strengthening mechanisms observed in the sample's tensile test. Since this XRD pattern demonstrates that the Al₂O₃ nanoparticles were not broken in any manner, they are discrete, non-deformable crystals; then any moving dislocations cannot cut through them. The dislocations should instead cut around the nanoparticles through the Orowan bowing mechanism [34]. The large temperature difference in the CTE between the hard ceramic particles and the compliant aluminum matrix leads to a small plastic zone that spreads around each nanoparticle during cooling after fabrication. This forms an excessive concentration of GNDs. It is this combination of harsh Orowan pinning and dense GND networks that causes the tremendous increase in the macroscopic yield stress, up to and including about 100 MPa.

The XRD data indicate that there are two metallurgically quite distinct hardening processes; both the Al-5%Mg solid-solution structure is characterized by a high level of atomic lattice distortion: the integration of the peaks and their broadening, and the uninterrupted formation of the reactive, new intermetallic compound(s) such as Al₃Mg₂. In the 5%-Al₂O₃ nanocomposite, the pattern shows only primary aluminum peaks and clear, consistent ceramic local maxima, with no new chemical reaction products. This crystallographic information ultimately demonstrates that the Mg alloy is strengthened by chemical reactivity and internal lattice strain. In contrast, the Al₂O₃ nanocomposite relies entirely on the physical and mechanical constraints of individual, non-deformable ceramic nanoparticles. Table 6 presents a comprehensive metallurgical comparison of the XRD findings for both synthesized metal-matrix systems.

Table 6. Comparative X-ray diffraction (XRD) analysis of AA1100-Mg and AA1100-Al₂O₃ composite systems

4.6 Analysis of the tensile test results

The microstructural analyses in the above sections, including those based on optical microscopy, SEM, TEM, and EDS, have identified the underlying crystallographic and phase changes resulting from the addition of Mg powder and Al₂O₃ nanoparticles to the commercially pure AA1100 supporting matrix. At this point, as shown, the characterization shifts from the crystallized, nanoscale level to the macroscopic mechanical response of these engineered systems; tensile testing is used as the critical diagnostic tool to confirm the structure-property relationships theorized in previous sections. The engineering stress-strain curves described below are used to quantitatively assess the effects of the various strengthening mechanisms on the kinematics of material deformation. In particular, this comparison is between solution-solidification hardening by the substitution of Mg with the atomic component, precipitation hardening, and the inflexible mechanical obstructions, Orowan bowing, and thermal dislocation networks induced by the non-deformable ceramic nanoparticles. Results after this point outline the traditional metallurgical trade-offs between improved yield strength and long-term tensile ductility, providing a conclusive assessment of the mechanical soundness of these two specific reinforcement approaches [35]. The distance between the macroscopic strain and yield behaviours of the reactive solid-solution system and the thermodynamically stable nanocomposite is large, so their respective stress-strain behaviours are plotted on separate axes below so as to maintain readability of the localized yielding behaviours.

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Figure 7. (a) The tensile test results for the AA1100 aluminum alloy with 5% Mg powder addition, and (b) the tensile test results for the AA1100 aluminum alloy with 5% Al₂O₃ nanoparticle addition

Figure 7(a) shows the results of the analysis of the stress-strain curve of the AA1100 matrix with the addition of 5% Mg powder, leading to the appearance of particular and successful micro-integration of the alloying element. A metallurgist views the curve as a recommended example of solid-solution strengthening and extraordinary strain-hardening ability. The curve shows a linear elastic region, followed by a shift in yield stress, after which it becomes macroscopically plastic at 40 ± 2 MPa. Substitutional solute atoms are incorporated in the FCC lattice of aluminum through the introduction of 5% Mg. Since the atomic radius of Mg (~1.60 A°) is much greater than that of aluminum (~1.43 A°), the atomic-scale compressive strain in the matrix is well-localized and spherical in nature [36]. These localized strain fields interact elastically with the stress fields of existing mobile dislocations. The Mg atoms pin the dislocations. A higher applied external stress is required to separate mobile dislocations within these solute atmospheres and to cause the material to exceed its elastic limit, resulting in an increased initial yield point compared with that of the base AA1100. Meanwhile, the most visible part of this curve is the gigantic homogeneous mass of plastic strains. The metal steadily increases in weight up to the yield point, then reaches the final tensile strength (UTS) of 146.5 ± 3.5 MPa, with an engineering strain of 60 ± 2.5%. Dislocation-dislocation interaction is the dominant strengthening mechanism across the vast domain and is essentially altered by the addition of Mg. The pure aluminum matrix has its stacking fault energy reduced by Mg. When the SFE is low, the partial dislocations have a greater distance between them, so the screw dislocations will find it much harder to cross-slip and to circumvent the obstacles. As a result, the dislocations are pushed aside, causing the total dislocation density to soar rapidly in the forest. Dynamic recovery would seem to be the limiting factor, which could have caused the large, constant strain-hardening rate. (σ/ϵ) observed on the curve. This is the very ability to continuously work-harden that enables the material to spread stress evenly and not localize into thinning to such an extent.

The strength of the material peaks near 145 MPa, which is much higher than that of commercially pure AA1100 in the annealed state (typically around 90 MPa). Since the Mg powder additions are complete, this curve is of outstanding secondary use for assessing the quality of material processing. In the case of poor powder herding, or in the presence of intense oxidation steps on previous particle domains, the material would undergo brittle breakdown during pre-testing or separation across granules, well before it could encounter such tension and strain. The continuous, uniform rise in UTS to a high level suggests very good solid-state alloying, very good inter- and intra-particle metallurgical bonding, and very good load transfer through the matrix [37]. Geometric instability (necking) does not occur until a strain threshold of about 60%. After the acute peak of tensile static deformation, the stress decreases rapidly; according to the considered criterion, macroscopic necking starts at an equivalent point where the material's rate of strain hardening is less than the actual stress (dσ/dε ≤ σ). Increasing the Mg content to 5% completely prevents this instability, owing to its ability to sustain the high strain-hardening rate. The drop-off in the direct vicinity, immediately after the UTS, is quite steep and hints that when localized void nucleation commences, which it will do at the location of any remaining intermetallic phases, microscopic pores, or the smallest undissolved Mg/oxide particles, void coalescence and ductile rupture take place rather rapidly.

An examination of this stress-strain curve in Figure 7(b) illustrates the deformation kinematics of an MMC in a textbook, incorporating 5 wt.% Al₂O₃ nanoparticles in an AA1100 matrix, essentially transforming the material from a high-ductility, low-strength metal to a high-strength structural composite.

The curve has a sharp initial elastic slope that reaches a yield stress of about 100 ± 4 MPa. It is an enormous improvement over unreinforced AA1100, which typically yields at approximately 30–40 MPa. The rigid Al₂O₃ nanoparticles, as opposed to the Mg addendum (which strengthens upon submersion in the atomic lattice), commonly enhance the matrix due to macro- and microscale mechanical constraints. In elastic deformation, stress is transferred from the object to the soft, compliant aluminum matrix via shear stress across the interface between the matrix and the high-modulus Al₂O₃ nanoparticles. As the fabrication cools, the large CTE mismatch between the ceramic particles and the aluminum matrix causes a localized plastic zone to form around each nanoparticle. This gives a high concentration of GNDs before testing the material. To initiate macroscopic yielding, the applied stress must overcome the resistance of the initial dislocation network.

At the very start of plastic flow, before smooth strain hardening sets in, a characteristic local perturbation, or "juggernaut," appears in the curve. This is one of the most important aspects to consider when writing, because anomalies that cannot be explained are easily flagged by reviewers. In MMCs supported by nanoparticles, this jagged yielding is usually attributed to a local accumulation of micro-stresses that suddenly bursts. Mobile dislocations then interact with the tightly packed ceramic nanoparticles as the matrix begins to flow plastically. It is thought that the serrations indicate the sequential jump of dislocations off pinning sites, or moments of microcracking in micro-sized agglomerates of nanoparticles, which leads to momentary stress reduction in advance of local work hardening of the surrounding matrix, followed by the re-establishment of flow stability.

After the yield aberration, the material undergoes vigorous, continuous strain hardening, steadily increasing to a final tensile strength of 212 ± 6 MPa at an elongation of 15 ± 1.2%. Due to the lack of coherence and shear ability of Al₂O₃ nanoparticles, moving dislocations cannot slice through them. Instead, they need to pass the nanoparticles around via the Orowan bowing mechanism. When dislocations are bowed around the nanoparticles, some of them will be left behind in the form of dislocation loops (Orowan loops). These loops represent new and extremely powerful barriers to later dislocations. This causes a sudden hardening of these forests, which exponentially increases the stress required to bring about further plastic deformation. The high, slanting, stainless-steel-like slope is uninterrupted graphical evidence of a high level of dislocation multiplication and entanglement, such that the matrix ends all primary plastic-strain accommodation at about 15% core strain.

Table 7. Mechanical response and deformation kinematics comparison

The observed decrease in the matrix's plastic strain potential can be correlated with the addition of 5 wt.% non-deformable ceramic particles. Under severe stresses, the large dislocation pile-ups at the interfaces between the matrix and particles give rise to tremendous local stress concentrations. Since the Al₂O₃ particles will not deform, the strain is relieved either through interfacial decohesion (the tearing away of the particle) or particle fracture [38]. These serve as nucleation points for the voids. The coalescence of voids occurs very rapidly once initiated, leading to a sudden macroscopic fracture at 15% strain. It is noted that a 15% elongation in a 5% nano-ceramic MMC is a strong outcome, implying high powder dispersion and strong interfacial bonding. A quantitative comparison of the macroscopic mechanical responses in the solid-solution Al-5% Mg alloy and the dispersion-strengthened Al-5%Al₂O₃ nanocomposite is presented in Table 7. It compares the high strain-hardening capability of the Mg addition with the significant enhancement in yield strength achieved by the ceramic nanoparticles.

4.7 Dynamic impact toughness and fractography

To put the static tensile data into perspective with the high-strain-rate application, the dynamic energy absorption of the composite systems was assessed using the Charpy V-notch impact test, as stipulated in ASTM E23. A solid-solution-enhanced system (AA1100 + 5 wt.% Mg) was shown to have outstanding cold energy-absorbing properties, with an average impact toughness of 59 ± 3 Joules. This macroscopic high-energy value strongly supports the observation that, under quasistatic tensile loading, there was extensive uniform plastic deformation (~60% elongation). The fact that the material could be loaded suddenly and violently to 59 Joules implies that the kinetic energy of the pendulum was distributed to fail the material through the rapid multiplication of dislocations and extensive plastic flow, and not through simple premature brittle cleavage at the hot end. To understand the specific micro-mechanisms that preserve this dynamic failure, the post-impact fracture surfaces were examined using SEM. In the fractography of the Al-5% Mg in Figure 8(a), the fracture morphology is classic and highly localized mixed-mode fracture with the overwhelming predominance of microvoid coalescence. Deep, equiaxed dimples that predominate on the fracture surface indicate ductile behavior. These physical residues of highly localized tearing of plastics are fibrous craters. Within the millisecond of contact, the extremely plastic aluminum matrix experienced extreme plastic flow, with the pendulum energy absorbed as the material physically tore while settling under the applied strain. But it is the secondary micro-constituents that determine when this fracture initiates.

The addition of 5 wt.% Al₂O₃ nanoparticles to the AA1100 matrix fundamentally transforms the characteristics of the high-ductility, low-strength metal into a high-strength structural composite, thereby reducing the impact energy.

The essential reduction in the matrix's plastic strain potential is due to increasing the percentage of non-deformable ceramic particles to 5%. Since the material is incapable of large-scale, uniform plastic deformation, it absorbs very little of the pendulum's kinetic energy before disintegration. The visual confirmation of the brittleness of the failure is the graphical depiction in Figure 8(b), which shows very irregular, blocky, and jagged topography dominated by the ceramic secondary phase. These sharp edges are likely to act as significant stress concentrations in an MMC. The material in the Charpy test experiences severe stresses during the violent impact, resulting in large numbers of dislocations at the interfaces between the material and the ceramic particles. It is this dynamic that creates extremely high levels of localized stress within the material.

Within some of the larger dimples are separate, terraced, blocky crystalline structures. These are the brittle intermetallic phases identified during synthesis, namely the equilibrium phases Al₃Mg₂ and Mg₂Si. These stiff intermetallics cannot plastically deform under heavy triaxial stress before the crack tip moves. So, they break through transgranular cleavage planes, becoming the dominant stress-concentration units and sites of imminent inner-hole formation. The discrete localized voids in the nucleated shattered intermetallic were then broken by the ductile matrix (or the α-Al) surrounding the dislocations, which rapidly strained itself to connect the voids and ultimately seal them [39]. The final result, which is 59 Joules of absorbed impact energy, is a direct measure of the mechanical work required to propagate the large-scale microvoid coalescence that establishes a highly ductile failure mode despite the alloy's substantially high yield strength. Adding nanoparticles usually increases the material's brittleness. The absorbed energy is 12 ± 1.5 Joules in the Al₂O₃ sample, indicating that under dynamic, high-strain-rate loading, the macroscopic failure is highly brittle.

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Figure 8. (a) AA1100 with 5% Mg powder, Charpy test sample fracture Scanning Electron Microscopy (SEM) morphology, and (b) AA1100 with 5% Al₂O₃ nanoparticles, Charpy test sample fracture SEM morphology

The SEM image depicts the minute micro-mechanisms of failure resulting from the lack of thermodynamic and mechanical compatibility between the reinforcement and the matrix. Since the highly fragile Al₂O₃ particles would not deform, they would be completely incapable of dissipating the strain imposed abruptly [39]. Therefore, the stress has to be relieved either by physical separation between the particle and the matrix, known as interfacial decohesion, or by the fracture itself of the ceramic particle. The craters are very irregular, and their sides are very sharp, as seen in the SEM image, providing physical evidence that these particles were crushed or violently ripped from the matrix. Locations of decohesion and fractured particles are instantaneously the nucleation sites for voids in the microstructure. Once voids are initiated at the damaged interfaces, they coalesce very quickly. This rapid interconnection of micro-voids prevents the surrounding aluminum from absorbing energy via progressive tearing, ultimately leading to an abrupt macroscopic fracture. This is broadly consistent with the summarized failure mode of this composite, which indicates a sudden macroscopic failure due to interfacial decohesion or rigid-particle fracture. The impact testing of Charpy demonstrates that the solid-solution Al-5% Mg alloy (59 Joules) possesses a high level of dynamic toughness as compared to the Al-5%Al₂O₃ nanocomposite (highly brittle) (12 Joules). Kinetic energy dissipation occurs efficiently in the Mg-alloyed matrix through extensive plastic flow and ductile microvoid coalescence, even before ultimate failure. Table 8 compares the dynamic impact toughness and fracture morphology results.

Table 8. Comparative analysis of dynamic impact toughness and fracture morphology for Mg and Al₂O₃ reinforced AA1100

Comparative evaluations of the tensile and dynamic impact data presented in the previous sections suggest the development of two distinct selection logics for applying these two reinforcement strategies. The in-situ solid-solution system (AA1100 + 5 wt.% Mg) showed outstanding values for dynamic energy absorption (59 ± 3 Joules) as well as high ductility (60 ± 2.5% elongation). This processing route is therefore strongly recommended for applications where impact resistance is desired, and structures are required to withstand loads without failure (e.g., automotive crumple zones or crash barriers, dynamically loaded housings, etc.). In contrast, the nanocomposite system (AA1100 + 5 wt.% Al₂O₃) exhibited exceptional tensile properties (UTS: 212 ± 6 MPa and yield strength: 100 ± 4 MPa) due to the strong Orowan pinning effect. This makes the Al₂O₃ route more appropriate for non-working applications where high stiffness and dimensional stability are prioritized, such as static load-bearing applications that are adequately protected against severe dynamic loading. However, it should be emphasized that the guidelines for material selection recommended herein are valid only within the set of investigated parameters: the reinforcement concentration was set at 5 wt.% only and the operational temperature was room temperature. Further experimental validation of extrapolating these guidelines to other volume fractions or higher temperatures is required.