On the spark plasma sintering of Mo-and-Cu containing low-density stainless steel: influence of sintering parameters on microstructure, densification and hardness considerations

Commercially available elemental powders of Fe, Mn, Al, Cr, C, Mo, and Cu, all with 99% purity procured from Sigma-Aldrich^®, were utilised. The powder mixing compositions listed in Table 1 were designed using a Processing–Input–Property (PIP) predictive modelling framework based on previously investigated composition datasets, with the aim of achieving duplex austenitic microstructures in alloys 1 and 3 and fully austenitic microstructures in alloys 2 and 4. Fully austenitic microstructures were targeted due to their high ductility, excellent corrosion resistance, and good biocompatibility, while duplex microstructures were designed to achieve a balanced combination of higher strength from ferrite and corrosion resistance and toughness from austenite. Mechanical alloying was employed to achieve potential homogeneity among the powders, which was conducted using a vertical ball mill. The parameters for mechanical alloying included a ball mill speed of 200 rpm, a ball-to-powder ratio of 20:1, a mixing time of 10 min, a relaxation time of 20 min, and a total effective mixing time of 12 h, culminating in a total mixing time of 36 h under inert gas. These parameters were adopted to prevent cold welding of the powder. The mixing was dry, and no process control agent (PCA) was used. Although process control agents are useful for limiting cold welding and particle agglomeration during mechanical alloying, they can also cause issues such as contamination, lower alloying efficiency, suppression of important microstructural developments, and the formation of unwanted phases [21‐23]. For applications where maintaining material purity or specific properties is essential, mechanical alloying is therefore sometimes performed without PCAs [21‐23]. Scanning electron images of the powders before and after mixing, are shown in Figs. 1 and 2.

The mass of the alloy to be sintered was calculated using Eqs. 1 and 2 where \(\:{w}_{i}\), \(\:{\rho\:}_{i}\), \(\:{m}_{alloy}\), \(\:r,\:h\) and \(\:\rho\:\) is the mass fraction of the elements, density of the alloy, mass of alloy, radius height and density respectively [24, 25]. The measured powders were sintered at different parameters, using temperature range between 900\(\:℃\) and 1100\(\:℃\), pressure of 40 MPa and 50Mpa, and hold time was kept constant at 5 min. The two sintering temperatures (900 °C and 1100 °C) were selected based on a previous optimisation study. The lower temperature of 900 °C was chosen for alloys 3 and 4 to prevent melting of the Cu additions, which could occur at higher temperatures, while still achieving adequate densification. The higher temperature of 1100 °C was applied to alloys 1 and 2 to evaluate maximum densification.

The sintered density and porosity were calculated using Eqs. 3 and 4 following Archimedes’ principle and ASTM B311 where \(\:{\rho\:}_{s}\), \(\:{m}_{d}\), \(\:{m}_{w},{m}_{s}\) and \(\:{\rho\:}_{o}\) is the density of the alloy, dry mass, wet mass, suspended mass and open porosity respectively [26‐28]. The results are reported in Table 2.

Samples were prepared by mechanically grinding using silicon carbide papers, followed by polishing with alumina suspensions to achieve a smooth, reflective surface suitable for SEM analysis.

The samples were subjected to hardness testing in accordance with ASTM E92 [29]. The samples were subjected to 5 different loads, 0.3kgf, 0.5kgf, 1kgf, 3kgf and 5kgf using the ALS Vickers hardness tester. The data obtained was used to calculate the hardness parameters using the Nix and Gao model [30]. The average depths (d) and indentation depth (h) was calculated using Eqs. 5 and 6 where D1 and D2 are the measured depths from the indenter.

Using the calculated data, a plot of \(\:{H}^{2}\) (hardness) against \(\:\frac{1}{h}\) is generated. The slope and intercept from this plot were then applied to Eq. 7 to determine \(\:{H}_{0}\)and \(\:{h}^{*}\), macroscopic hardness and characteristic length respectively.

Equation 8 through 13 were used to calculate the statistically stored dislocations (\(\:{\rho\:}_{SSD}\)), average shear strain (\(\:\gamma\:\)), geometrically necessary dislocations (\(\:{\rho\:}_{GND}\)), dislocation spacing, and total dislocation density (\(\:{\rho\:}_{T}\)) where \(\:\alpha\:,\mu\:\), \(\:b\), \(\:{L}_{SSD}\) and \(\:{L}_{GND}\) is the Taylor factor, shear modulus, Burger’s vector, statistically stored dislocations spacing and geometrically necessary dislocations respectively [30].

Examining powder morphology is essential for understanding how the shape and physical characteristics of powder particles influence the final properties of components produced through powder metallurgy. Morphology refers to the characterisation of powders’ physical form and particle geometry [31]. Although this study does not extensively explore the effect of particle morphology on sintering, the fundamental shape and structure of the powders were analysed to evaluate their potential for sintering. It is well established that particle size primarily affects powder compactability or compressibility, while particle shape significantly influences sinterability [31‐33].

Figure 1; Table 3 show the shape of the particles for iron (Fe), manganese (Mn), aluminium (Al), chromium (Cr), molybdenum (Mo) and copper (Cu). Molybdenum is spherical but agglomerated with smooth appearance. Manganese shows an angular (flaky) shape with sharp edges and a rough appearance. Iron shows a fibrous shape with a slightly rough appearance. Copper exhibits a dendritic shape with a smooth appearance. Chromium shows an angular shape with sharp edges and a rough appearance. Aluminium has a spherical shape with a rough appearance.

The particles used to produce the alloy exhibited a more angular morphology rather than a spherical one. This characteristic promoted rapid sintering over simple compression, leading to increased porosity at the expense of overall density [34, 35]. Given that the intended application of the alloy is biomedical implants, which typically require a porosity range of 10–30%, the rapid sintering behaviour is advantageous [36]. (Figure 2.)

The phases formed during mechanical alloying were identified using XRD and analysed with Panalytical Xpert HighScore software to evaluate the homogeneity of the mixed powders. The corresponding results are shown in Fig. 3; Table 4. Multiple diffraction peaks and varying cell volumes were observed, indicating the formation of several phases during mechanical alloying. The slight shift in the diffraction peaks between the different milled powders is attributed to composition-dependent changes in lattice parameters caused by substitutional effects or solid solution formation. Since all powders were milled under identical conditions, the peak shift reflects variations in the crystal lattice rather than milling parameters. This behaviour is consistent with Bragg’s law and Vegard’s law, as reported in mechanically alloyed powders [37]. For all four alloy compositions studied, the crystal structure was primarily cubic, exhibiting phase group numbers [229], [225], and [217], with dominant phases identified as Fe₂, Mo, α-Mn, and Al₄. The XRD patterns revealed a predominant BCC ferritic phase typical of Fe–Cr–rich alloys, along with FCC austenitic peaks [225] and α-Mn of primitive cubic structure [217]. These results indicate a dual-phase duplex structure composed of BCC and FCC phases, with possible Mo phase. This observation will be further confirmed through EDS mapping and BSE imaging of the sintered alloys.

Sintering was conducted using predetermined optimum sintering parameters established from a preliminary study. Sintering initiated at 400 °C for all sintering experiments, this temperature was selected to allow for enough thermal energy to be supplied to the atoms to migrate so that mass transport and interparticle is promoted. Sintering relies atomic diffusion to bond particles, enhance densification and increase mechanical strength [38]. At lower temperatures atomic mobility is negligible, making diffusion-driven sintering ineffective. The driving force for sintering is the reduction of surface energy, which requires adequate thermal activation. Shrinkage and densification typically begin between 350 °C and 500 °C, marking the effective temperature range for sintering, as noted by Xiong et al. [38].

Figure 4 represent the results for alloys 1–4 illustrating the relationships between sintering temperature, time, and shrinkage. During the initial 0–4 min, the sintering temperature remained stable, and minimal shrinkage was observed. This steady stage allows for precursor decomposition, particle contact, and early microstructural rearrangement, all crucial for achieving a uniform transformation before rapid densification occurs at higher temperatures [39]. This trend carries across all the alloy tested. Between 4 and 10 min, rapid shrinkage was observed across all alloys within the 400–1100 °C range. This is the main densification stage, driven by plasma effects and electrical discharges that enhance surface diffusion and accelerate neck formation between particles [39]. As densification proceeds, pores close, diffusion slows, and grain growth stabilises, marking the final densification stage [40]. Beyond approximately 10 min, the curves plateau, indicating that densification has reached completion with minimal further shrinkage.

The alloys in the same set, i.e. alloy 2 A and 2B, alloy 3 A and 3B, used the same set of sintering parameters, as such the relationship between sintering time and temperature is the same for each set. Alloys A and B were included for each alloy, except Alloy 1, to evaluate the reproducibility of the results and ensure that the observed microstructural and property trends were consistent. Sintering began at the same time, at the same temperature and plateaued at the same points and completed at the same time and temperature. The same can be said for the relationship between sintering time and shrinkage, there are slight differences in the shrinkage most notable in alloy 4. Alloy 4 A has shrinkage slightly below 4.5 mm while 4B has shrinkage slightly above 4.5 mm, this can be attributed to the morphology of the mechanically alloyed alloy powders. The element powders used were established to have large differences in particle sizes, this led to segregation during milling and non-uniform alloy powders. As such some of the alloys show slightly different sintering behaviour, shrinkage. The observed shrinkage behaviour in Alloys 3 and 4 cannot be attributed to particle size alone, as alloying elements such as Cu and Mo also influence sintering: small additions of elements like Mo have been shown to impede densification due to slower diffusion rates, whereas Cu‑rich phases can enhance atom mobility and densification pathways under SPS conditions, thereby affecting shrinkage kinetics beyond morphological effects [41, 42].

Figure 5 illustrates the relationships between temperature and shrinkage and temperature and densification rate. Across all alloys, shrinkage increased with rising temperature. However, the shrinkage rate curves displayed double peaks, indicating instability in the densification process. This effect was most prominent in alloys 1 and 2, suggesting nonuniform densification [43]. The impact of this instability is further evaluated through microstructural and corrosion analyses. Based on the trends observed for the alloys, it cannot be concluded which element has the favourable sintering behaviour on the relationships between time, temperature, shrinkage and shrinkage rate alone.

The best alloy of each composition were plotted together in Fig. 6, alloys 1 and 2 had the same sintering behaviour due to their similar sintering parameters, the same applies to alloys 3 and 4. Alloys 3 and 4 completed sintering faster compared to alloys 1 and 2, as they were sintered at a lower temperature (900 °C) and pressure (40 MPa). The sintering parameters affect the sintering duration, at elevated temperatures, thermal diffusion increases, generally promoting faster sintering and densification. However, the overall densification rate also depends on the diffusion kinetics between different elements, which can remain slow even under high temperatures, thus requiring extended sintering durations for complete densification [20, 44, 45]. In systems containing elements like Cr and W, the inherently sluggish diffusion between particles further necessitates prolonged sintering times, even at elevated temperatures and pressures [46].

The sintering conditions: temperature, pressure, and hold time; were optimized in a prior study, where their effects were discussed in detail. Despite the optimisation in the preliminary experiment, slight changes in the composition would affect the sintering parameters hence the results obtained in Table 5. The resulting sintered alloys exhibited relative densities ranging from 91% to 98%, with alloy 2 A achieving the highest density, as summarised in Table 5. Although each set of alloys (2 A and 2B, 3 A and 3B, and 4 A and 4B) were processed under identical sintering conditions per alloy set as shown in Table 2, slight differences in relative densities; particularly between alloys 2 and 4; were observed. These variations will be further examined through microstructural and phase analysis.

The sintered alloys in this study showed porosity values between 2.38% and 5.77%, with alloys of lower relative density exhibiting higher porosity, which, while below the ideal biomedical range, is still acceptable for certain implant applications. Porosity plays a critical role in bone tissue integration, as it facilitates cell migration, vascularization, and mechanical interlocking with surrounding tissue, thereby improving osseointegration and implant stability. The mechanical performance, specifically hardness and dislocation behaviour, was analysed using the Nix and Gao [30] model and strain gradient plasticity approach, with results presented in Fig. 6; Table 6.

Density and porosity influenced hardness; however, porosity had a more significant effect. Even minor increases in porosity can lead to substantial reductions in hardness [30]. This trend is evident in the results of alloy 2 A, which had the highest density and one of the lowest porosities, also exhibited the highest hardness and the lowest sensitivity to the indentation size effect. The difference between the densities and porosities of alloy 2 A and 2B is small, 98% and 97%, and 2.38 and 2.67% respectively, the difference in hardness these alloys is quite significant. A general correlation can be seen, higher density often aligns with higher hardness, however, among these factors porosity remains the dominant influence. For instance, alloy 4B had the third highest density among the tested sintered alloys, yet due to its highest porosity level, it recorded the lowest hardness. Porosity plays a crucial role in determining material hardness, with higher porosity consistently leading to lower hardness across metals, ceramics, composites, and coatings [47‐49].

Further insights into deformation behaviour were derived from the Nix and Gao plots. Figure 6a and b shows the plot for plot of hardness (H) versus indentation depth (h) and H² versus 1/h respectively. Figure 6a shows that hardness decreases with increasing indentation depth. At shallow depths, hardness is higher due to the formation of geometrically necessary dislocations (GNDs) that accommodate large strain gradients beneath the indenter, increasing resistance to deformation. As indentation depth increases, the strain gradient and GND density decrease, causing hardness to approach a constant bulk value (H₀). Thus, the H–h curve typically starts high at low depths and gradually levels off, reflecting the transition from GND-dominated deformation to bulk plastic behaviour. Figure 6b shows H² increases as 1/h increase. This indicates that more geometrically necessary dislocations (GNDs) form as the indentation depth decreases. This reflects a higher strain gradient and suggests that the material is strain hardening rapidly [30, 50].

Alloys 2 A and 2B, have higher hardness values in both Fig. 6a and b, this means these alloys are harder than alloys 1 A, 3 A and 4B across the range of depths measured. There is greater resistance to plastic deformation that can be attributed to its higher density as well as porosity. This means that alloy 2 has larger contribution to GNDs under different indentation meaning it effectively resists localized deformation. In contrast, alloys 1 A, 3 A, and 4B exhibited relatively flat and lower slopes, suggesting consistent deformation behaviour across different depths. These observations are supported by the dislocation data shown in Fig. 6d. Alloy 2 A had the highest GND density (Table 4); necessary to accommodate indentation; alongside one of the lowest h* values, suggesting it is closer to bulk-like deformation [30]. It also exhibited the highest statistically stored dislocation (SSD) density, indicating inherent material strength [51]. A direct comparison between Fig. 6b and d shows that steeper slopes in the hardness-depth plots correspond to higher total dislocation densities, further confirming the relationship between indentation behaviour and microstructural mechanisms.

A strong positive correlation exists between hardness and dislocation density, as increased dislocation density restricts dislocation movement, enhancing resistance to plastic deformation and resulting in higher hardness values [52‐55]. This trend can be observed in Fig. 6c and d, with the exception of alloy 2B, hardness increases from alloy 4B, 3 A, 1 A to 2 A and this correlates the total dislocation density increasing from 2 A, 1 A, 3 A to 4B. Factors such as grain size, alloy composition, and temperature can further influence the degree to which dislocation density affects hardness [54]. Alloy 2B has high hardness with lower total dislocation density, this suggests that mechanisms other than dislocation strengthening are dominant in resisting deformation [56, 57]. This can be attributed to a microstructure that is not uniform due to non-unform alloying powder. The microstructure instability compensates for the lower dislocation density, resulting in a material that is hard but not heavily strain-hardened [56, 57].

It is important to distinguish between the experimental hardness (H) and the true hardness (H₀), which reflects the intrinsic strength of the material without indentation size effects (ISE) [30, 50]. Due to ISE, H is often higher at shallow depths [58]. As shown in Fig. 6e, H₀ values are consistently lower than H, especially in porous alloys. This indicates that ISE must be considered when comparing hardness, as it can exaggerate strength if not accounted for.

The SEM-BSE images and EDS mapping results for alloys being studied are presented in Figs. 7, 8, 9 and 10. The microstructure plays a crucial role as it directly affects the hardness of the alloy [59]. The microstructures of alloys 1 (Fe-30.9Mn-4.9Al-4.5Cr-0.4 C) and 2 (Fe-21.3Mn-7.6Al-4.3Cr-1 C) appear to be similar due to their comparable compositions, the same applies to alloys 3 and 4. Mechanical alloying and XRD analysis of the powders showed the alloy powders had different particle size that resulted in non-uniform mixing and affected the packing state and void ratio of the final alloy. The EDS maps obtained from SEM suggests that the alloys are not homogeneous. Figures 7, 8, 9 and 10, show localised Cr-enriched, Al enriched, and Mo-enriched (alloy 3 and 4) areas characterised by the different intensities of colour. This uneven distribution impacts the overall microstructure. Although the composition was tailored to produce microstructures that are austenitic for alloys 2 and 4 and austenite-based duplex for alloys 1 and 3, the final microstructure for all alloys is austenite-based duplex. The deviation from the fully austenite microstructure can be attributed to mechanical alloying and sintering that introduced non-equilibrium effects such as carbide formation.

The Fe-Mn-Al-Cr-C alloy is a complex and composition-sensitive system, with the amounts of Al and Mn significantly influencing the balance between austenite and ferrite phases [60, 61]. In the initial mechanically alloyed powders, alloys 3 (Fe-30.9Mn-4.9Al-4.5Cr-0.4 C-3Mo-3Cu) and 4 (Fe-21.3Mn-7.6Al-4.3Cr-1 C-3Mo-3Cu) exhibited a dual-phase structure of austenite with Mo precipitates, while alloys 1 and 2 showed similar microstructures due to their comparable compositions. Alloys 1 and 2 primarily consist of γ-austenite with a minor α-ferrite (Fe-Mn) phase, as shown in Figs. 11. EDS mapping of these alloys indicates that Mn largely dissolves into Fe, consistent with its high solubility in iron and its role as an austenite stabiliser, often forming isomorphous carbides. Al, on the other hand, partially dissolves in Fe and partially forms intermetallic compounds, reflecting its limited solubility in iron; Al addition also promotes ferrite (BCC) stabilization [62]. Cr behaves similarly, partially dissolving into Fe and forming Cr-rich precipitates. Consequently, the microstructure comprises a γ-austenite matrix with embedded α-ferrite and Cr-rich precipitates, typical of many low-density steels with higher Mn and C content.

Alloys 3 and 4, containing Mo and Cu in addition to the base elements, display a similar dual-phase microstructure of γ-austenite and α-ferrite with precipitates and carbides. Copper acts as an austenite stabiliser and, due to the high sintering temperature, dissolves into the Fe-Mn-Al-Cr phase without forming precipitates, despite its low solubility in ferrite below 500 °C. Molybdenum, a ferrite stabiliser like Cr and Al, enriches the ferritic phase and can also form precipitates during cooling after austenitic transformation [62]. Overall, the microstructure of all four alloys is characterized by a γ-austenite matrix with an α-ferrite phase containing Cr- and Mo-rich precipitates, with additional stabilization effects from Mn and Cu.

The interrelationships among powder characteristics, sintering behaviour, resulting microstructures, and mechanical response of the studied low‑density stainless steels reveal the complex nature of processing–structure–property linkages in mechanically alloyed and spark plasma sintered (SPS) systems. Powder morphology and particle size have a fundamental influence on subsequent densification behaviour because they determine initial packing and the driving force for atom transport during sintering. Smaller particles with high surface area‑to‑volume ratios promote a greater driving force for diffusion, while particle shape governs contact area, packing efficiency, and ultimately the ease with which particle rearrangement and neck formation occur during SPS [63]. Spherical powders tend to pack more uniformly, facilitating predictable shrinkage and homogeneous density distribution, whereas irregularly shaped particles, with their additional contact points and higher surface roughness, can accelerate early interparticle bonding at the expense of final density and contribute to increased localised porosity [63]. Although spherical particles are generally favoured for loose powder sintering due to their superior rearrangement behaviour, irregular particles can enhance green strength through mechanical interlocking but often lead to higher porosity and limited densification in the final compact.

Mechanical alloying itself alters powder characteristics beyond morphology and size distribution. The high energy input associated with prolonged milling introduces lattice distortion, generates defects, and enhances diffusion kinetics to a degree not attainable by conventional processing. As a result, mechanically alloyed powders exhibit finer microstructures and increased chemical activity, enabling the formation of supersaturated solid solutions, metastable phases, or even amorphous structures depending on the element combination and milling parameters [14, 64].

Despite employing optimised sintering parameters, including temperature, pressure, and hold time determined from preliminary studies, small differences in alloy chemistry altered densification outcomes. Relative densities ranged from approximately 91% to 98% with corresponding porosity levels between 2.38% and 5.77%. Alloy 2 A attained the highest relative density, while alloys with nominally identical compositions (for example 2 A/2B and 4 A/4B pairs) exhibited measurable variation in density and porosity despite identical thermal cycles. This underscores the sensitivity of SPS densification to initial powder condition and localised compositional heterogeneity arising from mechanical alloying. The inverse relationship between density and porosity aligns with expectations: as densification proceeds and pores collapse, density increases and porosity decreases.

Porosity not only influences densification but also strongly affects subsequent mechanical behaviour. In the context of biomedical materials, porosity < 10% generally correlates with high mechanical strength but may limit biological integration due to reduced interconnectivity for tissue ingrowth [65, 66]. Porosity ranges of 5–25% are often dominated by closed pores, which impede vascularisation, while natural cortical bone exhibits porosity of 5–15%, creating a narrow window where mechanical support and biological functionality must be balanced [67‐69]. Although the sintered alloys in this study do not fall within the typical range idealised for maximal osseointegration, their lower porosity enhances mechanical stability, indicating potential suitability for load‑bearing implant applications where structural support is primary.

Hardness measurements interpreted via the Nix–Gao indentation size effect (ISE) framework further illustrate the combined influence of microstructure and porosity on mechanical response. The Nix–Gao model explains the increase in measured hardness at shallow indentation depths by the formation of geometrically necessary dislocations (GNDs) required to accommodate high strain gradients beneath the indenter [70, 71]. In this work, hardness consistently decreased with increasing indentation depth, reflecting a transition from strain–gradient dominated resistance to bulk plastic behaviour. True hardness values (H₀) derived from Nix–Gao extrapolation were consistently lower than measured values (H), particularly in alloys with higher porosity, emphasising the need to account for ISE when comparing mechanical response across specimens.

Among the alloy series, porosity emerged as a dominant factor over density in determining indentation hardness: alloys with lower porosity generally exhibited higher resistance to plastic deformation. For example, Alloy 2 A, with the lowest porosity values, recorded the highest hardness and lowest sensitivity to ISE, highlighting the central role of pore fraction in controlling local deformation behaviour. Conversely, Alloy 4B, despite relatively high bulk density, showed the lowest hardness due to its higher porosity. This observation is consistent with reported trends in metals, ceramics, and composites where hardness decreases either linearly or exponentially as porosity increases, as pores act as stress concentrators and collapse under applied load, reducing the effective load‑bearing cross‑Sects [30, 47‐49, 72‐74].

Further insight is gained from hardness versus indentation depth (H–h) and H² vs. 1/h plots, which reflect differences in GND accumulation among the alloys. Alloys 2 A and 2B exhibited steeper slopes, indicating rapid strain hardening and elevated GND densities. These trends are supported by total dislocation density measurements, where higher dislocation densities correlated with increased hardness, consistent with the positive relationship between dislocation barriers to plastic flow and measured hardness [52‐55]. Alloy 2B, while exhibiting comparatively high hardness, showed lower total dislocation density than expected; this suggests that mechanisms other than dislocation strengthening, such as microstructural heterogeneity or localised segregation, contributed to resistance against plastic deformation consistent with previous microscopy-based studies using combined SEM–EDS analysis that have demonstrated that such compositional and microstructural heterogeneity produces localized deformation behaviour, leading to significant differences in indentation response even within the same specimen [65‐67]. The hardness scatter observed in the present study can be attributed to microstructural non-uniformity inherent to powder-based spark plasma sintering. Similar behaviour has been widely reported in SPS alloys, where rapid heating rates and short dwell times can result in incomplete diffusion homogenization, localized elemental segregation, and heterogeneous phase distribution [65‐67]. These studies further show that hardness does not always correlate directly with global strengthening indicators, such as total dislocation density, because local compositional fluctuations and phase variations can dominate mechanical response [65]. Therefore, the hardness variability and deviation from expected strengthening trends observed in this work are consistent with established findings in SPS-processed alloys and primarily reflect spatial microstructural heterogeneity rather than intrinsic differences in overall strengthening mechanisms.

SEM–BSE imaging and EDS mapping revealed heterogeneous elemental distributions within the sintered microstructures. Although the compositions were designed to yield either fully austenitic or duplex microstructures, all alloys ultimately exhibited duplex γ‑austenite and α‑ferrite structures with Cr‑ and Mo‑rich precipitates. This deviation from the intended fully austenitic structure likely arises from non‑equilibrium effects associated with mechanical alloying and the rapid thermal cycles of SPS, which can stabilise ferritic phases and carbides despite equilibrium predictions. The Fe–Mn–Al–Cr–C system is particularly composition‑sensitive, with Mn stabilising the γ‑austenite phase and Al and Cr promoting α‑ferrite, while Cu, as an austenite stabiliser, tends to fully dissolve into the matrix at SPS temperatures and Mo enriches ferritic regions and forms Mo‑rich precipitates upon cooling [60‐62]. The presence of dual phases, compositional segregation, and precipitates affected localised strain gradients beneath the indenter and contributed to the observed ISE behaviour.

Taken together, these findings illustrate that the mechanical performance of SPS low‑density alloys is governed by an interplay of processes: mechanical alloying establishes initial microstructure and compositional uniformity, sintering densification and pore evolution dictate residual porosity and density, and both microstructural heterogeneity and dislocation mechanisms govern resistance to deformation. Alloys with higher density, lower porosity, and more uniform microstructures exhibited superior hardness and reduced ISE sensitivity. While this study provides a comprehensive exploration of processing–structure–property trends under fixed SPS conditions, future work involving systematic parameter variation and comparative benchmarking will be required to fully quantify the relative influence of individual variables and to generalise these insights for broader alloy design contexts.