Processing, Microstructure, and Mechanical Behavior of Tungsten Heavy Alloys for Kinetic Energy Penetrators: A Critical Review

Abstract

Tungsten heavy alloys (WHAs) are two-phase composites known for their exceptional density, strength, hardness, and ductility, making them ideal for radiation shielding, kinetic energy penetrators, and aerospace components. Due to their high melting point, WHAs are primarily processed via powder metallurgy, with liquid-phase sintering (LPS). Spark plasma sintering (SPS) and microwave sintering are emerging as advanced consolidation techniques. Recent research has focused on improving WHA performance through microstructural manipulation, alloying with elements like Fe, Co, Mo, and Re; rare earth oxides like Y₂O₃, La₂O₃, and Ce₂O₃; and employing high-entropy alloys (HEAs) as matrix phase. Additionally, additive manufacturing (AM) techniques are increasingly being used to fabricate complex WHA components. Despite their advantages, WHAs still exhibit limitations in penetration performance, primarily due to their tendency to form mushroom-like heads upon impact rather than self-sharpening. Ongoing research seeks to enhance shear localization, refine grain structure, and optimize processing methods to improve the mechanical properties and impact resistance of WHAs. Furthermore, modeling and simulation approaches are being explored to understand the mechanical behavior of WHAs. This review comprehensively overviews the above aspects and presents recent advances in WHA processing.

1. Introduction

Kinetic energy penetrators (KEPs) are high-velocity projectiles designed to defeat armored targets using their mass and speed for penetration [1]. Modern KEPs, particularly Armor-Piercing Fin-Stabilized Discarding Sabot (APFSDS) rounds, require materials with high density and strength to maximize their impact effectiveness [2]. An ideal KEP material should exhibit low work-hardening rates and strain-rate hardening to facilitate flow softening and promote shear localization at lower strains. It should also have a low heat capacity to allow for a rapid temperature increase, along with high density and strength to maximize impact energy. Additionally, it should possess a high thermal-softening rate to ensure adiabatic shear banding initiation at lower temperatures, thereby enhancing penetration performance [3]. The performance of KEPs depends on various factors, including the penetrator’s material properties, geometry, mass, velocity, penetration depth, failure mechanisms, and target characteristics.

WHAs are two-phase composites made of W particles embedded in a ductile metallic matrix, typically containing 85–95 wt.% W [3]. W is a refractory metal with a melting point of 3420 °C [4], requiring extremely high temperatures for consolidation, and sintered pure W parts tend to be brittle [5]. To address these challenges, W powders are mixed with lower melting metals or metal combinations with some solubility in W [6]. The WHA matrix is a solid solution primarily of Ni, often with additives such as Fe, Co, Cu, Re, Ta, and Mo [6]. During sintering, the low-melting elements form a liquid phase, allowing rapid consolidation into a fully dense two-phase structure. Advanced sintering techniques, including SPS and microwave sintering, enable improved densification at lower temperatures with faster processing times [7].

WHAs are widely used for KEPs due to their high density (17–19 g/cm³), exceptional strength, and good ductility. They also offer reasonable toughness, radiation shielding, moderate electrical and thermal conductivity, effective vibration damping, and strong resistance to corrosion. Beyond KEP applications, WHAs are utilized in various industrial and military components, including high-performance weapon systems, rocket nozzles, oilfield rejuvenation projectiles, golf club weights, self-winding watch weights, aviation wing weights, and industrial vibrators [6,7,8,9,10]. Their high density enhances kinetic energy retention, leading to deeper armor penetration, while their toughness and resistance to deformation ensure effectiveness upon impact [11]. Importantly, WHAs serve as a viable alternative to depleted uranium (DU) KEPs, offering a non-radioactive and environmentally safer option [6].

The primary drawback of WHAs is their lower penetration capability compared to DU alloys [6]. The superior penetration ability of DU alloys is attributed to their self-sharpening behavior, which arises due to the formation of adiabatic shear bands (ASBs) upon impact. These bands promote localized shear failure, allowing the penetrator to maintain a sharp profile, thereby enhancing penetration depth [12]. In contrast, WHAs tend to form a mushroom-like tip, which reduces their overall penetration efficiency. Since ASBs play a critical role in penetration performance, key research focus for WHAs has been focused on enhancing their adiabatic shear localization tendency. This has been pursued through compositional modifications, such as adding Co [13] and Mo [14] to the Ni matrix or by replacing the Ni matrix with elements like Hf [15,16], Ti [17], and Ni₃Al [18,19,20]. Additionally, novel processing approaches have been explored to enhance the mechanical properties of WHAs [13,21].

This review article summarizes recent advancements in WHA processing, strategies for producing high-strength WHA through microstructural manipulation, matrix composition modifications to promote ASB formation, thermo-mechanical processing techniques, and computational approaches for WHAs.

2. Processing of WHAs

WHAs are generally processed either by liquid-phase sintering (LPS) [22,23,24] or solid-state sintering (SSS) [25,26]. As previously discussed, WHAs consist of two components: the W phase and the matrix phase. The primary difference between LPS and SSS lies in the sintering temperature. In SSS, the sintering temperature remains below the melting point of both components, keeping the matrix phase in a solid state with no liquid formation. In contrast, LPS involves heating above the matrix phase melting point, allowing it to liquefy and bind the W particles.

Various techniques have been explored for sintering WHAs, including conventional high-temperature furnaces, SPS, and microwave sintering. These methods differ mainly in heating mechanisms (radiation/convection, microwave heating, and resistance heating) and heating rates. Conventional furnace sintering typically achieves heating rates of around 3–20 °C/min [27], while microwave sintering achieves rates around 50 °C/min [28], and SPS exceeds 100 °C/min [29]. Rapid heating and shorter dwell times in SPS and microwave sintering result in refined microstructures with enhanced mechanical properties [30]. Infiltration techniques have also been employed for WHA processing [31,32]. For instance, W-Cu [31] and W-Cu-Ag [33] alloys are fabricated through infiltration techniques, where a molten matrix phase permeates a sintered W skeleton [10]. Compared to conventional sintering, infiltration produces finer microstructures but increases W-W particle interfaces, leading to higher contiguity and reduced ductility.

While conventional sintering methods provide dense parts, they limit the fabrication of complex geometries. Powder injection molding (PIM) addresses this by enabling intricate shapes at lower machining costs [34,35,36]. In PIM, WHA powders are mixed with a polymer or wax binder to create a feedstock, which is then injection-molded into the desired shape [37,38]. The part then undergoes debinding to remove the binder, followed by sintering to achieve high density and desirable mechanical properties [39]. However, PIM faces challenges such as complex debinding procedures, dimensional instability, and microstructural defects, requiring careful optimization. Zu et al. [35] optimized the injection-molding of 90W-4.9Ni-2.1Fe, achieving a 1100 MPa tensile strength and 25% elongation, with properties influenced by the solvent-de-bonding temperature and environment [6]. AM techniques such as selective laser melting (SLM) [40,41,42,43], direct energy deposition (DED) [44,45,46], binder jet printing (BJP) [47], and bound metal deposition (BMD) [48,49] offer new possibilities for fabricating complex WHA components with minimal material waste and greater design flexibility. However, WHA processing via AM presents challenges due to W’s high melting point, low laser absorptivity, high thermal conductivity, and brittleness, leading to defects such as microcracking and poor densification [50].

Irrespective of the processing route, WHAs often exhibit poor mechanical properties due to the presence of a significantly weaker matrix phase than the W phase, brittle W/W interfaces. Defects such as porosity further degrade performance. To enhance strength and hardness, post-sintering thermo-mechanical treatments, such as hot isostatic pressing (HIP) [51], forging [52], rolling [53], extrusion [54], and swaging [22,55,56] have been employed. The following sections provide a comprehensive review of WHA processing techniques, highlighting their advantages, limitations, and potential for further development.

2.1. Sintering of WHAs

Extensive research has been conducted on the sintering of WHAs with diverse alloy compositions such as W-Ni-Cu [57,58,59], W-Ni-Fe [26,27,60,61,62,63,64,65,66,67,68,69], W-Ni-Fe-Co [13,27,56,58,70,71,72], W-Ni-Fe-Co-Mo [14,73,74,75,76,77,78], W-Ni-Co [79,80,81,82], W-Ni-Cu-Fe [14,83], W-Ni-Fe-Re [21,67], W-Ni-Fe-Mo [14], W-Ni-Mn [84], and W-Ni-Cu-Mn [85]. Recently, bulk metallic glass (BMG) W-ZrCuAlNi [86,87] and HEAs have been explored as alternative matrix materials, such as W-AlCrFeNiV [88], W-CoCrFeMnNi [89], W-FeNiCoCrCu [90], and W-CoCuFeNi [91].

Sintering is a thermal process that densifies powder compacts at elevated temperatures below their melting points. Since WHAs consist of a W phase and a matrix phase, the elemental powders must first be mixed to ensure uniformity. Various mixing techniques, including turbula mixing, attritor milling, planetary ball milling, high-energy vibratory milling, and mechanical alloying, are employed to achieve homogeneous powder blends [26,61,81,92]. Ball milling can function as either a high-energy mechanical mixing process or a low-energy blending technique. In high-energy milling, the powder particles undergo repeated cold welding, rupturing, and re-welding, altering their structure [93]. In contrast, low-energy milling improves homogeneity and particle bonding through controlled collision and deformation. Tumbler mixing, a low-energy approach (with or without milling media), ensures uniform elemental distribution without significant particle size reduction or structural changes.

Selecting the appropriate mixing method is essential for achieving high-density WHA samples. For 90W-7Ni-3Fe and 92.5W-5.25Ni-2.25Fe compositions, attritor milling enhances tensile properties and ductility over turbula mixing due to the improved powder dispersion [94]. Chen and Ma [95] investigated two milling methods for a W-Ni-Co-Y₂O₃ alloy. In the first, all elemental powders were ball milled for 24 h. In the second method, the Ni and Co powders were pre-milled separately for 24 h before being combined with W powder and Y₂O₃ nano powder for an additional 24 h. The latter approach produced a more homogeneously distributed matrix phase, resulting in improved microstructural uniformity.

The next step after mixing is compaction to the required shape. Compaction minimizes porosity, enhances densification, and ensures structural uniformity. Before compaction, WHA powders are typically reduced at 700 °C in a H₂ atmosphere for 2 h to remove the oxide layer that develops during the mixing process [56,70]. This reduction step improves sintering behavior, removes impurities, and relieves stress from the mixing process. Panigrahi et al. [90] reported that mechanically milled W and HEA powders benefit from heat treatment at 600 °C for 2 h in H₂ to relieve cold working stresses induced by mechanical milling. Generally, WHA powders are compacted using cold isostatic pressing (CIP) or hydraulic pressing. CIP applies uniform pressure in all directions, yielding higher green and sintered densities compared to uniaxial hydraulic pressing [96]. Table 1 summarizes the different WHA powder mixing and compaction methods.

Once compacted, WHAs undergo sintering, where process parameters, such as temperature, time, and atmosphere, play a critical role in determining microstructure and mechanical properties. Optimal sintering conditions promote densification, fine W phase dispersion, uniform matrix distribution, and minimized W/W interfaces [97]. Generally, WHAs are sintered in a reducing H₂ atmosphere to prevent W oxidation at high temperatures. This effectively removes surface oxides, leading to cleaner metal surfaces and improved densification [14,22,70]. The W particle size and distribution also significantly affect sintered material properties. Fine W particles enhance densification and mechanical strength due to their larger surface-area-to-volume ratio, facilitating faster sintering [98]. However, finer W powders may also increase the risk of grain coarsening at higher sintering temperatures, which can reduce toughness. The following section presents the literature on the various sintering techniques used for WHA consolidation.

Table 1. Overview of the powder mixing and compaction methods used in the processing of WHAs.

Composition (wt.%)	Mixing Technique and Duration	Ball-to-Powder Ratio	Compaction Methods	Ref.
93W-4.9Ni-2.1Fe	Tumbler mixing for 30 h	-	CIP at 180 MPa	[60]
92.5W-6.4Ni-1.1Fe	Turbula mixing for 30 min	-	Hydraulic pressing at 200 MPa	[61]
90W-6Ni-2Fe-2Co	Ball milling for 24 h	2:1, SS balls	CIP at 200 MPa	[27]
93W-4.2Ni-1.2Fe-1.6Co	Ball milling for 48 h at 50 rpm	2:1	CIP at 250 MPa	[99]
90.5W-7.1Ni-1.65Fe-0.5Co-0.25Mo	Ball milling at 35 rpm	1:1, SS balls	CIP at 250 MPa	[77]
90 W-6Ni-4Cu	Turbula mixing for 10 h at 96 rpm	-	Hydraulic pressing at 400 MPa	[83]
91W-6Ni-3Co	Drum mixing for 20 h	-	Hydraulic pressing at 200 MPa	[82]
90W-5Ni-2Fe-3Co	Ball milling for 48 h at 40 rpm	1:1	CIP at 250 MPa	[70]
90W-7Ni-3Fe	Vibratory milling at 1000 rpm	-	Hydraulic press	[26]
95W-3.5Ni-1.5Fe	Planetary ball milling for 5 h at 180 rpm	5:1	CIP at 200 MPa	[92]

Table 1. Overview of the powder mixing and compaction methods used in the processing of WHAs.

Composition (wt.%)	Mixing Technique and Duration	Ball-to-Powder Ratio	Compaction Methods	Ref.
93W-4.9Ni-2.1Fe	Tumbler mixing for 30 h	-	CIP at 180 MPa	[60]
92.5W-6.4Ni-1.1Fe	Turbula mixing for 30 min	-	Hydraulic pressing at 200 MPa	[61]
90W-6Ni-2Fe-2Co	Ball milling for 24 h	2:1, SS balls	CIP at 200 MPa	[27]
93W-4.2Ni-1.2Fe-1.6Co	Ball milling for 48 h at 50 rpm	2:1	CIP at 250 MPa	[99]
90.5W-7.1Ni-1.65Fe-0.5Co-0.25Mo	Ball milling at 35 rpm	1:1, SS balls	CIP at 250 MPa	[77]
90 W-6Ni-4Cu	Turbula mixing for 10 h at 96 rpm	-	Hydraulic pressing at 400 MPa	[83]
91W-6Ni-3Co	Drum mixing for 20 h	-	Hydraulic pressing at 200 MPa	[82]
90W-5Ni-2Fe-3Co	Ball milling for 48 h at 40 rpm	1:1	CIP at 250 MPa	[70]
90W-7Ni-3Fe	Vibratory milling at 1000 rpm	-	Hydraulic press	[26]
95W-3.5Ni-1.5Fe	Planetary ball milling for 5 h at 180 rpm	5:1	CIP at 200 MPa	[92]

2.1.1. Solid-State Sintering of WHA

SSS of WHAs involves heating compacted powders below the melting point of both components (matrix and W). Since no liquid phase forms, densification is primarily driven by atomic diffusion, particle rearrangement, and grain boundary migration [100]. Additionally, plastic deformation mechanisms such as grain boundary sliding [101] and dislocation motion contribute to densification [102]. However, grain growth can occur as surface energy decreases, with larger grains forming at the expense of smaller ones, altering the material’s microstructure [102,103]. Ryu et al. [104] investigated the mechanical alloying of 93W-5.6Ni-1.4Fe and found that 48 h of milling produced a nanocrystalline W grain size of 16 nm. When sintered at 1300–1400 °C, the material achieved ultrafine W particles (~3 μm) and a high density (>97%). In contrast, liquid-phase sintering (LPS) at 1485 °C led to coarser particles (~27 μm) and a lower density due to swelling. Similarly, Gurwell [25] found that finer W powders (~1 μm) improved sinterability but reduced ductility, while coarser powders (5–10 μm) enhanced ductility but required longer sintering times. Optimal sintering near the matrix solidus temperature provides a balance between strength and ductility by incorporating both fine and coarse W powders.

Despite its advantages, SSS presents challenges. While SSS produces refined WHA microstructures, densification is slower due to its dependence on solid-state diffusion, often resulting in residual porosity. Additionally, maintaining a uniform distribution of the W and matrix phases is difficult due to atomic mobility and solubility differences, which can create weak regions and cause mechanical property variations. These limitations make SSS less widely used compared to LPS.

2.1.2. Liquid-Phase Sintering of WHAs

During LPS, the sintering temperature exceeds the liquidus temperature of the matrix phase. However, densification depends on selecting low-melting elements that can dissolve W, as only these elements effectively promote grain wetting and rearrangement. Pure Cu has a negligible solubility for W and fails to achieve significant densification, despite forming a liquid at sintering temperatures (Figure 1a). In contrast, adding Ni to Cu enhances wetting and enables partial W dissolution, resulting in near-complete densification [105]. As a result, early WHAs have Ni and Cu (90W-6Ni-4Cu and 90W-5Ni-5Cu, density ~16.5 g/cc) [106]. More recently, WHAs have evolved to include transition metals like Fe and Co, which offer superior solubility for W. Among these, W-Ni-Fe alloys have been extensively studied [105,106].

Unlike SSS, LPS significantly accelerates microstructural evolution by enhancing material transport through the liquid phase. Capillary forces pull W grains together, facilitating particle rearrangement [3]. Sharp W particle edges dissolve preferentially, leading to grain shape accommodation, while finer W grains dissolve into the molten matrix and later reprecipitate during cooling, a process governed by Ostwald ripening [3,107]. Sintering temperature and dwell time are critical in determining the final microstructure (Figure 1b). Suboptimal conditions result in excessive/insufficient liquid formation, causing inhomogeneity. Kuncicka et al. [81] investigated the effect of sintering temperature on 93W-6Ni-1Co and found that sintering below 1525 °C for 30 min produced an inhomogeneous structure. In contrast, sintering at or above 1525 °C or increasing the dwell time to 60 min led to improved homogeneity, enhanced mechanical properties, and greater W diffusion.

Higher W content increases the W volume fraction, contiguity, and grain size. Contiguity, which measures interphase contact, below 0.38 has been shown to prevent shape distortion, as confirmed by percolation theory and experimental findings [108]. Heating rates commonly used during the LPS of WHAs typically range from 5 to 10 °C/min. Bollina and German [109] analyzed densification at heating rates of 1, 5, 10, and 15 °C/min for WHAs with 83–93 wt.% W in a Ni-Fe (7:3) matrix. They found that the highest sintered density (95%) was achieved at 1 °C/min for 83 wt.% W.

Despite achieving near-theoretical densities [3], LPS WHAs are prone to shape distortion or slumping due to gravity [110]. This occurs because of density differences between solid W and liquid phases, leading to W particle segregation. LPS is typically restricted to materials with a high solid-phase fraction, and shorter sintering times are preferred [107].

Table 2. Summary of sintering conditions used for LPS of WHAs.

Composition (wt.%)	Sintering Cycle and Sintering Atmosphere	Ref.
93W-4.9Ni-2.1Fe, 93W-4.2Ni-1.2Fe-1.6Co, 93W-4.9Ni-1.9Fe-0.2Re	Sintering temperature: 1480 °C Holding time: 2 h in H2 Heating rate: 8 °C/min Cooling rate: 1 °C/min	[99]
89.75W-6Ni-2Fe-2Co-0.25Mo	Pre-sintering temperature: 1300 °C/1 h Final sintering temperature: 1480 °C Holding time: 75 min in H2 Heating rate: 3–5 °C/min	[78]
94.9W-3.4Ni-1.7Cu	Sintered temperature: 1500 °C Holding time: 1.5 h in H2	[14]
90W-5.7Ni-3.8Cu-0.5Fe	Degassing temperature: 1000 °C/20 min Final sintering temperature: 1400 °C Holding time: 1 h in H2	[83]
93W-6Ni-1Co	Sintering temperature: 1525 °C Holding time: 30 min in H2	[81]
88W-xNi-yCu; (x = 8.94–9.60, y = 2.4–3.06)	Debinding temperature: 500 °C/1 h/5 °C/min Pre-sintering: 500–1000 °C/1 h/5 °C/min/AC Final sintering temperature: 1412 °C Holding time: 1 to 10 min in H2.	[57]
93W-4.9Ni-2.1Fe	Sintering temperature: 1493 °C Holding time: 3 h	[63]
91W-6Ni-3Co	Sintering temperature: 1520 °C Holding time: 20 min in H2	[82]
90W-7Ni-3Fe	Pre-sintering: 1350 °C/15 °C/min/H2/15 min Then, till 1350–1450 °C/10 °C/min Final sintering: 1510 °C/5 °C/min/vacuum	[69]
88W-8.4Ni-3.6Fe	Sintering: 1500 °C/30 min in H2/10 °C/min up to 800 °C, and 5 °C/min thereafter. During cooling, atmosphere changes from H2 to Ar.	[66]
93W-4.9Ni-2.1Fe	Pre-sintering: 1380 °C/2 h under H2/AC, then microwave sintering: 1480 °C/10 min/10 at.% H2 + 90 at.% N2/15 °C/min	[64]
97W-2.1Ni-0.9Fe, 99W-0.7Ni-0.3Fe	Sintering: 1600 °C/3h/Ar-20 vol.% and H2-200 mL/min	[111]
89W-7Ni-3Fe-1Re	Sintering temperature: 1480 °C Holding time: 2 h in H2	[21]
92W-7.2Ni-0.8Co, 92W-6.4Ni-1.6Co, 92W-5.6Ni-2.4Co, 92W-4.8Ni-3.2Co, 92W-4Ni-4Co	Sintering temperature: 1520 °C Holding: 20 min in H2	[80]
90W-6Ni-4Co, 90W-7Ni-3Co, 90W-8Ni-2Co	Pre-sintering temperature: 1300 °C/1 h Final sintering: Between 1480 and 1490 °C Holding time: 75 min in H2	[79]
94.9W-3.4Ni-1.7Cu, 92.5W-5Ni-2.5Cu, 96.1–2.8Ni-1.1Cu, 96W-2.8Ni-1.1Cu-0.1Fe, 93.5W-4Ni-2Fe-0.5Co, 90.5W-7.2Ni-1.8Fe-0.45Co-0.05Mo, 74W-8Ni-2Fe-16Mo	Sintering temperature: 1500 °C Holding time: 2 h in H2	[14]

summarizes the sintering cycles and atmospheres used during the LPS of various WHA systems.

2.1.3. Spark Plasma Sintering of WHAs

SPS, also known as pulsed electric current sintering, is an advanced powder metallurgy technique that enables rapid densification through the simultaneous application of high temperature, pressure, and electric current [29]. In SPS, a graphite die containing the powder mixture is subjected to 30–70 MPa pressure, while an electric current generates uniform heating via resistance, enabling rapid densification at lower temperatures and shorter holding times [112]. Unlike LPS using a conventional furnace, which requires extended holding times, SPS minimizes the sintering time, limiting grain coarsening and leading to finer microstructures and improved mechanical properties. Hu et al. [113] investigated W-Ni-Fe alloys processed via SPS and found that slower heating rates promoted better densification and finer grain size, while higher rates accelerated transient liquid-phase formation, causing uneven microstructures and porosity. Pan et al. [114] studied mechanically milled 90W-6Ni-4Mn powders sintered at 1150 °C, achieving a high density (99.6%) and superior strength due to uniform matrix phase distribution and controlled grain growth. However, higher sintering temperatures increased porosity and reduced strength.

Using one-time spark plasma sintering on 90W-6Ni-4Mn powders (Figure 2a), W grain size was refined below 10 µm, achieving a maximum hardness of 69.1 HRA at 1150 °C [115]. Fan et al. [29] applied two-time spark plasma sintering (TTSPS) to the same alloy (Figure 2b), initially sintering at 1000 °C, followed by a second sintering at 1100–1200 °C with heating rates of 100 °C/min and 200 °C/min, respectively. This yielded finer W grains, improving hardness (76.6 HRA) and bending strength (785.3 MPa) due to the optimized grain structure and matrix phase distribution (Figure 2c). Singh et al. [116] analyzed an 84W-7Ni-3Fe-6Nb WHA sintered via SPS, observing that melt squeezing at 1215 °C indicated transient liquid-phase sintering due to eutectic reactions at powder contacts. The highest density (99%) was achieved at 1200 °C for 84W-7Ni-3Fe-6Nb, while 90W-7Ni-3Fe and 84W-7Ni-3Fe-6Mo reached 97.9% and 95.9% at 1150 °C, respectively. The 84W-7Ni-3Fe-6Nb alloy exhibited the highest hardness (678 HV at 1150 °C) due to the persistent liquid phase, whereas 84W-7Ni-3Fe-6Mo peaked at 589 HV at 1200 °C before declining. Coatings like boron nitride reduce carbon diffusion from graphite dies and prevent carbide formation on the sample surface (e.g., WC).

SPS enables rapid densification of WHAs at lower temperatures and shorter processing times (

Table 3. Summary of the literature data on SPS parameters used in the processing of WHAs.

Composition (wt.%)	Sintering Condition	Reference
93W-5.6Ni-1.4Fe	Sintering temperature: 1050–1410 °C Pressure: 30 MPa Heating rate: 45–270 °C/min	[113]
99W-1Ni-1Co-1Fe-1Mo-1Cu	Sintering temperature: 1200 °C Pressure: 20 MPa Heating rate: 1200 °C/min	[117]
90W-6Ni-4Mn	Sintering temperature: 1000 °C to 1200 °C Pressure: 30 MPa Heating rate: 100 °C/min	[114]
93W-5.6Ni-1.4Fe	Sintering temperature: 1360–1430 °C Pressure: 30 MPa Heating rate: 10–380 °C/min	[65]
95W-3.5Ni-1.5Fe	Sintering temperature: 900–1300 °C Pressure: 50–70 MPa Heating rate: 50–300 °C/min	[118]
90W-5.6Ni-2.4Fe-2Co	Sintering temperature: 1400 °C Pressure: 30 MPa Heating rate: 100 °C/min	[13]
90W-6Ni-4Mn	First-step solid-phase-sintered at 1000 °C/100 °C/min/6 min, second-step sintering: 1000–1200/200 °C/min/3 min	[29]
90W-7Ni-3Fe, 84W-7Ni3Fe-6Mo, 84W-7Ni-3Fe-6Nb	Sintering temperature: 1150–1275 °C Pressure: 30 MPa Heating rate: 100 °C/min	[116]
88W-2Re-7Ni-3Fe	Sintering temperature: 1100 °C Pressure: 50 MPa Heating rate: 100 °C/min	[67]

), improving density and mechanical properties. Controlled heating rates help prevent porosity and microstructural inconsistencies. However, SPS is mainly suitable for smaller components and is limited in producing large parts such as KEPs.

2.1.4. Microwave Sintering of WHAs

Microwave sintering is another advanced technique that consolidates materials using microwave radiation, in which heating occurs internally (inside-out), unlike conventional sintering, through dipolar polarization, conduction losses, and magnetic field interactions [119]. Microwaves consist of electric and magnetic components that influence the movement of dipoles, free electrons, domain walls, and electron spins during processing [119]. Solid metals and other conductors reflect microwaves, causing surface heating and plasma formation while microwave-absorbing materials convert radiation efficiently into heat [120]. Materials that directly absorb microwaves, such as ceramics and metallic powders, undergo direct heating. Bulk metals, which cannot absorb microwaves directly, require a susceptor, a microwave-absorbing material like graphite or silicon carbide (SiC), for hybrid heating [121]. Hybrid heating occurs in three stages: (1) The susceptor absorbs microwave energy and heats up, (2) it transfers heat via conduction and convection to the non-absorbing material, and (3) once the material reaches a critical temperature, it begins absorbing microwaves directly, leading to rapid heating [119]. Unlike conventional sintering, which transfers heat from the surface inward, microwave sintering generates heat volumetrically from the core outward, enhancing efficiency and uniformity [122]. Initially, heat transfer is driven by conduction and convection, with radiation becoming significant once the susceptor surpasses a threshold temperature.

Microwave sintering of WHAs has been carried out using a multi-mode cavity with a microwave frequency of 2.45 GHz [61,64,123]. SiC susceptors like dense SiC, oxide-bonded SiC, or nitride-bonded SiC are commonly used due to their high thermal stability (~1600 °C) and structural integrity, ensuring efficient and uniform heating [124]. Compared to conventional furnaces with slow heating rates (5 °C/min), microwave heating reduces the processing time and energy consumption by ~75% [125]. Zhou et al. [125] investigated the microstructural evolution of WHAs under different sintering conditions. Microwave sintering of 90W-6Ni-2Fe-2Co at 1470 °C for 60 min resulted in finer and more homogeneous microstructures compared to conventional sintering for 120 min at the same temperature, attributed to rapid densification and reduced grain growth (Figure 3a,b). Further, 90W-7Ni-3Fe alloys microwave-sintered at 1480 °C showed noticeable microstructural densification after just 10 min of holding time (Figure 3c,d). As shown in Figure 3e, increasing the heating rate promoted finer W grains and reduced contiguity, both of which are beneficial for mechanical performance.

Prabhu et al. [126] reported that microwave sintering significantly improved densification in pure W, achieving a 93% theoretical density with high-energy milled powder compared to 85% with as-received powder due to the finer particle size and higher surface energy. Microwave-sintered WHAs exhibit finer W particles and lower contiguity than conventionally sintered alloys, leading to superior mechanical properties [27]. For instance, a 90W-6Ni-2Fe-2Co alloy sintered via microwave reached a tensile strength of 989 MPa and impact strength of 203 J/cm², while its conventionally sintered counterpart achieved only 816 MPa and 17 J/cm². Microwave-sintered 90W-7Ni-3Cu samples also exhibited higher hardness than their conventionally sintered counterparts [59]. Liu et al. [127] studied the microstructural and mechanical behavior of a 93W-4.9Ni-2.1Fe alloy fabricated via microwave at 1500 °C. The process enhanced W dissolution and diffusion in the Ni-Fe matrix, yielding a relative density of 98.8%, tensile strength of 1185.6 MPa, elongation of 16.4%, and hardness of 42.1 HRC. As the sintering temperature increased, porosity decreased, W grains coarsened, and fracture modes shifted from intergranular to transgranular cleavage, improving mechanical strength. Additionally, W solubility in the matrix reached 35.79%, enhancing bonding strength while preventing brittle-phase formation due to rapid sintering and cooling.

Microwave sintering presents a cost-effective, energy-efficient alternative to conventional sintering, enabling faster heating, refined microstructures, and enhanced mechanical properties. However, MW heating develops non-uniform microstructures due to the skin effect, which causes surface-dominant heating and can lead to reduced electrical conductivity [7,128]. The skin effect occurs when alternating electromagnetic fields, like those used in microwave energy, induce electrical currents (called eddy currents) on the surface layer of conductive materials. These currents do not penetrate deeply into the bulk of the material, especially at high frequencies, and are confined to a shallow region near the surface, referred to as the “skin depth”. During microwave sintering, this concentration of current on the surface can lead to non-uniform heating, especially in materials like W, which is a good electrical conductor. While powdered metals can absorb microwave energy more effectively than bulk metals, surface-dominant heating still persists to some degree. As a result, there may be localized overheating and uneven microstructural evolution, which can introduce micro-defects, porosity, or grain boundary irregularities, all of which contribute to a drop in the material’s bulk electrical conductivity after sintering. Moreover, any phase transformations, impurity inclusion, or grain boundary oxidation that occur preferentially near the surface due to microwave-induced thermal gradients may further degrade the continuity of electron pathways, reducing overall conductivity. Strategies like using susceptors, hybrid sintering, and optimizing powder characteristics aim to improve heat distribution and minimize defects [7,128].

Table 4. Summary of the microwave sintering conditions used for the processing of WHAs.

Composition (wt.%)	Sintering Condition	Reference
92.5W-6.4Ni-1.1Fe	Sintering temperature: 1500 °C Holding time: 20 min in H2 Heating rate: 20 °C/min	[61]
90W-7Ni-3Fe	Sintering temperature: 1460 °C Holding time: 1 h in H2 Heating rate: 20 °C/min Susceptor: SiC	[27]
90W-7Ni-3Fe	Sintering temperature: 1480 °C Holding time: 10 min in H2 Heating rate: 10–112 °C/min Susceptor: SiC	[28]
90W-6Ni-2Fe-2Co	Sintering temperature: 1470 °C Holding time: 60 min in H2 Heating rate: 20 °C/min Susceptor: SiC	[27]
93W-4.9Ni-2.1Fe	Sintering temperature: 1480 °C Holding time: 10 min in H2 Heating rate: 15 °C/min	[64]
90W-7Ni-3Cu	Sintering temperature: 1300–1450 °C Heating rate: 22 °C/min	[59]
93W-4.9Ni-2.1Fe	Sintering temperature: 1510 °C and 1520 °C Holding time: 1 h in H2 Heating rate: 20 °C/min	[129]
93W-4.9Ni-2.1Fe	Pre-sintering temperature: 800 °C Sintering temperature: 1250–1500 °C Holding time: 5 min in 10 vol.% H2 and 90 vol.% N2 Heating rate: 30 °C/min	[127]
93W-4.9Ni-2.1Fe	Extruded-rod-sintered at 1550 °C Holding time: 30 min in 10 vol.% H2 and 90 vol.% N2	[130]
93W-4.9Ni-2.1Fe	Sintering temperature: 1250–1500 °C Holding time: 5 min in 10 vol.% H2 and 90 vol.% N2 Heating rate: 30	[131]

summarizes the research on WHA consolidation via microwave sintering.

A summary of the microstructural and property benefits from different sintering methods is presented in

Table 5. Comparative overview of microstructural features and mechanical property improvements associated with different sintering techniques used for WHAs.

Property	Microwave	SPS	Conventional
Process Time	Very Low	Very Low	High
Grain Size	Fine	Very Fine	Coarse
Density	>98.5%	95–98%	~98%
W Solubility in the Matrix	Low	Low	High
W-W Contiguity	Low-Moderate	Low	High
Matrix Uniformity	High	High	Moderate
Hardness (HV)	High	High	Moderate
Tensile/Impact Strength	Superior	High	Low
Distortion Risk	Low to Moderate	Low to Moderate	Moderate to High

.

2.2. Additive Manufacturing of WHAs

AM technologies have been an alternative for producing complex, customized WHA components, such as 90W-7Ni-3Fe [132,133,134], 95W-5Nb [40], 95W-5Ta [135], W-Cr [136], 75W-25Cu [137], W-2.5TiC [138], W-10 Cu [139], 90W-6Ni-2Fe-2Co [140], 80W-10Ni-10Cu [141], and pure W [142]. Powder bed fusion (PBF), a key AM method, uses either a laser (L-PBF) or electron beam (EBM) to selectively melt micron-sized powder layer by layer [143,144]. However, processing W is challenging due to its high ductile-to-brittle transition temperature [142]. Schwanekamp et al. [145] investigated the effect of L-PBF processing parameters on W-Ni-Fe alloys and found that optimal density was achieved at energy inputs of 500–1000 J/mm³, while excessive energy caused porosity due to evaporation. Preheating at 800 °C had minimal impact, but post-processing heat treatment at 1500 °C significantly improved the microstructure, making it comparable to conventionally sintered alloys. Similarly, Ivekovic et al. [146] reported that SLM-processed 90W-7Ni-3Fe reached ~95% density at 250–350 J/mm³, but excessive energy or preheating at 400 °C led to cracking. The as-built alloy exhibited brittle fracture with a UTS of 871 MPa.

In L-DED, WHAs such as W-Ni and W-Ni-Fe displayed layered microstructures with unmelted W and intermetallics, resulting in high strength but low ductility due to porosity and brittle fracture paths [147,148,149]. Wei et al. [150] studied the microstructure and fracture characteristics of a 90W-7Ni-3Fe alloy produced by L-DED. Compared to conventional LPS, L-DED resulted in a non-uniform microstructure with finer W particles and a higher W solubility in the matrix (Figure 4). Mechanical properties varied along the build direction, with the highest tensile strength in the middle section (Figure 4c). L-DED samples exhibited W particle cleavage and matrix rupture, while LPS samples showed more W-W interparticle fractures. Zhang et al. [132] reported that, during the laser sintering of W-Ni-Fe, the W particles did not fully melt or disperse evenly in the region with a cellular structure. In contrast, the Ni and Fe particles melted completely and fused with the unmelted W particles. EBM, with preheating up to 1400 °C, offered superior control over cracking and produced dense, crack-free structures with columnar grains and mixed textures [151,152,153,154,155]. Mechanical properties varied, with compressive strengths from 900 to 1523 MPa via alloying and oxide dispersion [151,152,153,154,155].

BJP is another viable AM method for producing WHA components with complex geometries. BJP offers fewer processing challenges than laser-based methods. The process involves inkjet printing a liquid binder onto a layer of metal powder to form a green part, which is then de-bonded and sintered to create a solid component [156]. Stawovy et al. [134] studied the BJP of W-Ni-Fe-Cu, achieving sintered densities of up to 99.7% of the theoretical value with 18–19% dimensional shrinkage. The tensile properties met ASTM standards, with an average UTS of 770 MPa and 8.6% elongation. Zhou et al. [136] developed a short-time laser powder sintering method that drastically reduced the liquid-phase duration, preventing excessive W grain growth and producing finer grains. This method resulted in WHAs with a 42% higher yield strength and improved ballistic performance. The challenges during AM can be mitigated by optimizing alloy compositions, refining powder selection, adjusting laser parameters, and improving processing environments.

Continued advancements in alloy design, microstructure control, and impurity mitigation are essential for producing high-performance WHAs via AM. Cracking is the most critical issue during the AM of WHAs due to W’s high ductile-to-brittle transition temperature (200–400 °C), thermal stresses, and grain boundary embrittlement from impurities such as oxygen [157,158,159]. Cracks predominantly form along high-angle grain boundaries (HAGBs) [157,158,160,161,162,163]. Although oxygen levels as low as 30 ppm have been reported, cracking persists, indicating that factors such as residual stress, grain morphology, and the intrinsic brittleness of W also contribute significantly [159,164]. Grain refinement, achieved through alloying additions (e.g., Ta, Nb, ZrC) or processing strategies (e.g., scan vector rotation, remelting), can reduce crack density but not fully eliminate it [159,165,166]. Powder characteristics (size, shape, distribution) and laser absorptivity also influence cracking behavior. W’s high thermal conductivity (173 W/m·K) accelerates heat dissipation and cooling, impairs melt pool wetting, and promotes balling defects [167]. Moreover, W’s high melting point, low laser absorptivity, and high viscosity hinder effective melting and fusion during printing. These challenges can be mitigated by using finer, spherical powders to enhance absorptivity and by optimizing process parameters to improve densification and microstructure control [168]. A systematic study on these parameters is still lacking.

Despite these inherent challenges, advances in laser processing, powder engineering, preheating, and hybrid manufacturing techniques have significantly improved build quality and mechanical performance. For example, Chen et al. [140] optimized SLM parameters (350 W laser power, 400 mm/s scan speed, 60 µm hatch spacing, 30 µm layer thickness) for W-18Ni-6Fe-6Co (wt.%), achieving 96.1% relative density, 1198 MPa ultimate tensile strength, and 9.5% elongation. Similarly, Wang et al. [147] optimized the process parameters for 90W-7Ni-3Fe (800 W laser power, 400 mm/min scan speed, 0.5 mm layer height, and a powder feed rate of 10–15 g/min under high-purity Ar shielding). The samples showed an alternating layered microstructure with reduced porosity (~0.87%), finer grains, increased W content in the matrix, and lower contiguity. These microstructural improvements, combined with supersaturated W in the matrix and enhanced interface bonding from rapid thermal cycling, yielded a high yield strength of 822 MPa and a tensile strength of 1037 MPa, surpassing conventional LPS WHAs. Hybrid AM approaches, such as combining AM with conventional techniques like sintering, hot isostatic pressing (HIP), or machining, are particularly beneficial for processing WHAs. One example is the BMD process, which integrates fused filament fabrication (FFF) and PIM to print bound metal rods into complex geometries, followed by debinding and sintering to produce dense metal parts. Bose et al. [48] used BMD to prepare 93 wt.% W, achieving a more than 96% sintered density.

2.3. Post-Processing of WHAs

For KEP applications, high strength and impact toughness are essential for optimal performance against target materials. However, WHAs, in their sintered state, often fall short of these requirements. To address this, post-sintering strategies such as heat treatment and thermo-mechanical processing (e.g., swaging and hot isostatic pressing) are employed to enhance mechanical properties. Various heat treatment methods have been explored, including vacuum annealing, solution treatment, cyclic heat treatment, and aging in different atmospheres (e.g., vacuum, H₂, N₂, Ar, or gas mixtures), followed by quenching in oil, water, or Ar. These treatments help prevent impurity segregation at the W/matrix interface and reduce hydrogen embrittlement, promoting a uniform distribution of elements within the matrix [58,169]. The heat treatment temperature is selected based on the solidus and solvus temperatures of the matrix.

For most WHAs, vacuum annealing at 1100–1250 °C followed by quenching can eliminate trapped hydrogen, remove W/matrix interfacial segregation, suppress intermetallic formation, and improve chemical homogeneity [76,79]. Das et al. [75] observed a decrease in W volume fraction (from 80% to 75%) and contiguity (from 0.7 to 0.3) in an LPS 90.5W-7.1Ni-1.65Fe-0.5Co-0.25Mo alloy after vacuum heat treatment at 1100 °C for 1 h. During this process, W partially dissolves into the matrix and remains in a solid solution after oil quenching. Solution heat treatment dissolves intermetallic compounds that may form during furnace cooling after sintering. Kiran et al. [73] found that in 90W-6Ni-2Fe-0.5Co-1.5Mo alloy, after three heat treatment cycles, contiguity and dihedral angle were reduced while matrix phase fraction was increased. This indication is relatively less prevalent in the case of single-cycle specimens.

Kumari et al. [71] reported that water quenching after solution treatment improved elongation and impact resistance in 89.6W-6.2Ni-1.8Fe-2.4Co alloys compared to Ar quenching. While Ar heat treatment removes hydrogen, water quenching prevents interstitial segregation and intermetallic precipitation at grain boundaries. During cyclic heat treatments, Kumari et al. [79] applied a repeated sequence of low-temperature holds (800–900 °C for 3 h) and high-temperature treatments (1100–1200 °C for 1.5 h), followed by vacuum heat treatment and water quenching. The low-temperature stage led to the formation of Co₃W intermetallic compounds, while the high-temperature stage dissolved them. More cycles increased the volume of intermetallic precipitates. Das et al. [58] reported improved tensile strength (from 717 MPa to 1000 MPa) and elongation (from 5% to 20%) in a 91W-7Ni-1.5Fe-0.5Co alloy after heat treatment due to improved W/matrix bonding and a more uniform impurity distribution.

Grain refinement through alloying and deformation via swaging can further improve WHA strength. Despite the challenges posed by WHA hardness and brittleness, advancements in swaging technologies have enhanced processing efficiency and final product quality. Swaging after sintering also increases density and promotes strain hardening. Kiran et al. [169] studied the effects of swaging and heat treatment on an LPS 90W-6Ni-2Fe-2Co alloy. Heat-treated rods were swaged and aged at various temperatures for 1 h in N₂. Swaging increased dislocation density, and subsequent aging significantly improved hardness. In a related study, Kiran et al. [73] found that cyclic heat treatment of 90W-6Ni-2Fe-0.5Co-1.5Mo reduced the contiguity and dihedral angle while increasing W solubility and the matrix volume fraction. Swaged specimens exhibited a lower work-hardening exponent than heat-treated ones. Using finite element analysis, Kocich et al. [30] used finite element analysis and experiments to assess sintered and cold-swaged W-Ni-Co alloys. Microhardness values for the sintered W particles and matrix were 389.2 and 296.8 HV, respectively. Swaging increased UTS significantly but reduced ductility. Caliskan et al. [69] investigated rotary swaging of a 90W-7Ni-3Fe alloy produced by vacuum LPS. They reported improved density, tensile strength, and ductility. Fractographic analysis revealed a shift from intergranular fracture in sintered alloys to a combination of transgranular fracture, matrix deformation, and W grain intergranular failure in swaged samples. The mechanical properties of WHAs after various post-processing treatments are summarized in

Table 6. Summary of mechanical properties of WHAs following various post-processing treatments.

Composition (wt.%)	Processing Condition	YS (MPa)	UTS (MPa)	Elong. (%)	Impact Strength (J/cm2)	Work-Hardening Exponent (n)	Ref.
90W-6Ni-4Co	Swaged (0.36 true strain) with six intermediate heat treatments (800–900 °C/3 h and 1100–1200 °C/4 h)	1334	1372	4	19	0.002	[79]
90W-7Ni-3Co	Swaged (0.36 true strain) with six intermediate heat treatments (800–900 °C/3 h and 1100–1200 °C/4 h)	1300	1339	10	281	0.003
90W-8Ni-2Co	Swaged (0.36 true strain) with six intermediate heat treatments (800–900 °C/3 h and 1100–1200 °C/4 h)	1236	1278	13	279	0.003
90W-7Ni-3Co	Swaged (0.36 true strain) with eight intermediate heat treatments (800–900 °C/3 h and 1100–1200 °C/4 h)	1341	1410	13	227	0.004
90W-8Ni-2Co	Swaged (0.36 true strain) with eight intermediate heat treatments (800–900 °C/3 h and 1100–1200 °C/4 h)	1247	1311	14	258	0.002
90W-6Ni-2Fe-2Co	Swaged (0.71 true strain) with two intermediate heat treatments (1100 °C/1.5 h)	1442	1461	5	71	0.05	[78]
89W-6Ni-2Fe-3Co	Swaged (0.71 true strain) with two intermediate heat treatments (1100 °C/1.5 h)	1540	1582	4	70	0.04
89.5W-6Ni-2Fe-2Co-0.5Mo	Swaged (0.71 true strain) with two intermediate heat treatments (1100 °C/1.5 h)	1522	1551	6	53	0.05
89.75W-6Ni-2Fe-2Co-0.25Mo	Swaged (0.71 true strain) with two intermediate heat treatments (1100 °C/1.5 h)	1522	1547	5	47	0.05
90W-6Ni-1.5Fe-2.5Co	Swaged (0.71 true strain) with two intermediate heat treatments (1100 °C/1.5 h)	1561	1594	5	73	0.04
90W-6Ni-1Fe-3Co	Swaged (0.71 true strain) with two intermediate heat treatments (1100 °C/1.5 h)	1575	1612	6	121	0.05
93W-4.9Ni-2.1Fe	Swaging to a 44% reduction with two intermediate heat treatments (1100 °C/2 h)	1255	1286	3	13	0.018	[99]
93W-4.2Ni-1.2Fe-1.6Co	Swaging to a 44% reduction with two intermediate heat treatments (1100 °C/2 h)	1219	1280	12	42	0.044
93W-4.9Ni-1.9Fe-0.2Re	Swaging to a 44% reduction with two intermediate heat treatments (1100 °C/2 h)	1347	1380	5	26	0.026
90W-7Ni-2Fe-1Co	Total of 41% deformation with one intermediate heat treatment (1150 °C/2 h)	1360	1400	10	65	-	[56]
93W-4.9Ni-1.4Fe-0.7Co	Total of 41% deformation with one intermediate heat treatment (1150 °C/2 h)	1410	1435	7	45	-
95W-3.5Ni-1.0Fe-0.5Co	Total of 41% deformation with one intermediate heat treatment (1150 °C/2 h)	1418	1420	4	14	-

.

3. Effect of Composition on the WHA Characteristics

To enhance the strength, hardness, and ductility of WHAs, strengthening the matrix phase through the addition of various alloying elements is one effective strategy. Common additions include Co, Mo, Ta, and Re. The W-Ni-Cu alloy was the first WHA developed. While it has a higher density than W-Ni-Fe alloys, its tensile strength is lower. Additionally, increasing Ni content in W-Ni-Cu alloys reduces the dihedral angle, further enhancing densification [24]. Das et al. [14] found that adding just 0.1 wt.% Fe to W-Ni-Cu alloys improved both the hardness and tensile strength. Shakunt et al. [83] reported that incorporating 1.0 wt.% Fe, while maintaining a Ni–Cu ratio of 6:4, increased the sintered density to 99.4%. However, a higher Fe content also increased the dihedral angle, indicating a more stable W/W interface than the W/liquid interface. This impeded liquid penetration and led to a higher compressive strength due to the reduced matrix volume and higher W/W contiguity. Wensheng et al. [170] observed that adding up to 1.0 wt.% Re to 93W-4.9Ni-2.1Fe increased the tensile strength from 997.2 MPa to 1161.2 MPa, though elongation dropped significantly from 26.4% to 8.6%. In contrast, Cr additions (0–1 wt.%) led to a reduction in tensile strength (from 997.2 MPa to 844.4 MPa) and elongation (from 26.4% to 7.7%). Cr tends to accumulate at interfaces, weakening interfacial cohesion and thereby degrading mechanical performance. Wang et al. [171] investigated W-Ta alloys with 5–20 wt.% Ta and found consistent grain size reduction compared to pure W. W-10Ta, prepared via SPS, demonstrated the best mechanical properties. Fractography revealed a transition from intergranular fracture in pure W to transgranular fracture in W-Ta alloys, indicating increased toughness with Ta addition. Chen et al. [172] reported that adding up to 10 wt.% Mo to W-Mo-Ni-Fe composites reduced the W grain size from 27.13 μm to 14.93 μm. Mo also improved the density and hardness, likely due to intermetallic compound formation.

Manganese (Mn) additions led to non-uniform elemental distribution in sintered compacts, reducing electrical conductivity, microhardness, and densification [85]. While traditional W-Ni-Fe WHAs perform well at low strain rates, their penetration capabilities are inferior to DU, likely due to a lower tendency to form ASBs. Since adiabatic shear is influenced by thermal conductivity and Mn has much lower thermal conductivity than Fe, replacing Fe with Mn may promote shear band formation. Hong et al. [173] studied a 90W-6Ni-4Mn alloy and emphasized the need to prevent manganese oxide formation during sintering to achieve full densification in Mn-containing WHAs.

The graph compiled from over 60 data points from the literature illustrates the effect of various alloying elements Fe, Cu, Mo, Re, Co, and Mn on the W content in the matrix phase of W-Ni-based WHAs (Figure 5). The summary of the effect of alloying elements on WNi-based WHAs is presented in Table 7. Among these, Co exhibits the highest W solubility, with both mean and median values clearly exceeding those of the other alloying additions. An increased W content in the matrix not only contributes to solid solution strengthening but also enhances the wetting and interfacial bonding between the matrix and W particles, which is critical for effective load transfer and improved performance under high-strain-rate conditions. In contrast, Cu shows the lowest matrix W content, reflecting its poor ability to dissolve W and limited wetting characteristics. This can result in weak particle/matrix interfaces and reduced mechanical properties. Mn shows a broad distribution in matrix W content, but its mean and median values are not significantly higher than those of the baseline W-Ni-Fe system. This suggests that, while Mn may influence local solubility or microstructural features, but due to its tendency to segregate, it does not substantially enhance overall W dissolution in the matrix. Re and Mo also exhibit moderate to low matrix W contents, with values close to or slightly below those of the Fe-containing baseline. Despite Re known advantages in high-temperature systems, its effect on W solubility in this context appears limited, possibly due to restricted solubility or kinetic factors during processing. Overall, matrix W content is a key factor influencing the strength, ductility, and interfacial integrity of WHAs. Among the elements studied, only Co clearly improves W solubility, making it particularly beneficial for high-performance applications such as KEPs.

The formation of fine W precipitates within the matrix of WHAs is governed by a complex interplay of alloy composition, thermal treatment, and microstructural evolution. Notably, alloys with a W-Ni-Co composition, such as 90W-7Ni-3Co, have shown a propensity to form fine W-rich precipitates within the matrix when subjected to specific cyclic heat treatments. These treatments typically involve an initial hold at 850 °C, where a W-rich intermetallic phase is formed in the Ni-Co matrix, followed by heating to 1150 °C to dissolve the intermetallic phase. Upon cooling, this process results in the nucleation of fine W precipitates (~1 μm), leading to a two-phase matrix microstructure that significantly enhances the alloy’s yield strength and strain-hardening response [174,175]. In contrast, WHAs containing Fe, such as 89.6W-6.2Ni-1.8Fe-2.4Co, tend not to form matrix precipitates under the same thermal conditions. This suppression of precipitation is attributed to the altered thermodynamic stability of the matrix phase caused by Fe addition, which discourages the formation and dissolution of the intermetallic phase necessary for precipitation [175]. Further studies have shown that the chemistry of the matrix itself plays a dominant role in determining precipitation behavior, where the W solubility and mobility in the matrix are key controlling factors [79]. Moreover, the presence of alloying elements like Mo has been shown to promote the precipitation of Mo-Ni intermetallic phases at W matrix interfaces when sufficient Mo is present in the liquid phase during sintering [172]. These findings underscore the importance of both thermodynamic driving forces and kinetic pathways in designing WHAs with tailored precipitation behavior and enhanced mechanical properties.

Additionally, rare earth oxides such as Y₂O₃, La₂O₃, and Ce₂O₃ have shown potential in refining the WHA microstructure and enhancing its mechanical performance, as fine oxide particles can improve WHA’s self-sharpening and penetration behavior. Dong et al. [176] synthesized W-5 wt.% Y₂O₃ composite powders via a wet chemical method, incorporating ultrasonic treatment and sodium dodecyl sulfate as a surfactant. They concluded that the traditional wet chemical process yielded high-quality composite precursors. Y₂O₃ additions improved hardness and thermal conductivity due to grain refinement. Park et al. [177] added 0.1 wt.% Y₂O₃ to a 93W-4.9Ni-2.1Fe powder mixture mechanically alloyed for 72 h, and LPS at 1493 °C for 1 h. Although oxide dispersion was achieved, the resulting specimens had a lower tensile strength and elongation than conventional LPS samples. Incomplete densification and micropores led to microcrack initiation under tensile loading, reducing ductility. These defects primarily formed at triple junctions and W/matrix interfaces. Xiao et al. [178] employed a hydrothermal method and mechanical alloying to fabricate nanoscale Zr(Y)O₂-dispersed 93W-4.9Ni-2.1Fe alloys using molecular-level liquid–liquid doping combined with hot isostatic pressing. The resulting alloy showed a more uniform nanoparticle distribution than conventional powder processing methods. This combined approach effectively reduced particle size, contributing to improved dispersion strengthening. The brittleness of WHAs is highly sensitive to microstructural factors such as W grain size, porosity, and impurity segregation at grain boundaries. To achieve fine-grained WHAs, ultrafine composite powder precursors and rapid sintering are effective.

Table 7. Summary of the effect of alloying elements on WNi-based WHAs [61,85,170,179,180].

Feature/Elements	Fe	Co	Mo	Re	Mn
Effect on the matrix	Solid solution strengthening	Solid solution strengthening	Solid solution strengthening	Strong solid solution strengthening	Limited solid solution strengthening
Grain size effect	Grain refinement at moderate Fe content; excessive Fe can lead to grain coarsening	Grain refinement at moderate Co content	Grain refinement and coarsening resistance	Strong grain refinement	Limited grain refinement
Sintering behavior	Improves wettability up to moderate Fe content; reduces at high Fe content	Improves wettability up to moderate Co content; reduces at high Fe content	Enhances densification	Enhances sintering kinetics and diffusivity	Improves wettability; assists in densification
Ductility	Decreases due to the formation of brittle intermetallic phases	Decreases above 6 wt.% due to brittle intermetallic phases	Maintains or slightly decreases	Improves both room- and high-temperature ductility	Slightly decreases
Intermetallic compounds	Fe7W6	Co3W, Co7W6	MoNi-type intermetallic compound	No intermetallics	Specific intermetallic compound not reported; however, MnO can form during sintering and can reduce density
Precipitate formation	Fe-rich precipitate forms at the W/matrix interface	Co-rich precipitates at the W/matrix interface	Mo-rich precipitates at grain boundaries	None, remains in solid solution	None
Specific mechanical properties	High strength and moderate ductility	High strength and moderate ductility	High-temperature strength and creep resistance	Superior high-temperature strength and ductility	Not ideal for high-temperature applications

Table 7. Summary of the effect of alloying elements on WNi-based WHAs [61,85,170,179,180].

Feature/Elements	Fe	Co	Mo	Re	Mn
Effect on the matrix	Solid solution strengthening	Solid solution strengthening	Solid solution strengthening	Strong solid solution strengthening	Limited solid solution strengthening
Grain size effect	Grain refinement at moderate Fe content; excessive Fe can lead to grain coarsening	Grain refinement at moderate Co content	Grain refinement and coarsening resistance	Strong grain refinement	Limited grain refinement
Sintering behavior	Improves wettability up to moderate Fe content; reduces at high Fe content	Improves wettability up to moderate Co content; reduces at high Fe content	Enhances densification	Enhances sintering kinetics and diffusivity	Improves wettability; assists in densification
Ductility	Decreases due to the formation of brittle intermetallic phases	Decreases above 6 wt.% due to brittle intermetallic phases	Maintains or slightly decreases	Improves both room- and high-temperature ductility	Slightly decreases
Intermetallic compounds	Fe7W6	Co3W, Co7W6	MoNi-type intermetallic compound	No intermetallics	Specific intermetallic compound not reported; however, MnO can form during sintering and can reduce density
Precipitate formation	Fe-rich precipitate forms at the W/matrix interface	Co-rich precipitates at the W/matrix interface	Mo-rich precipitates at grain boundaries	None, remains in solid solution	None
Specific mechanical properties	High strength and moderate ductility	High strength and moderate ductility	High-temperature strength and creep resistance	Superior high-temperature strength and ductility	Not ideal for high-temperature applications

4. Microstructure of WHAs

The microstructure of WHAs undergoes significant transformations during LPS, which directly impacts their mechanical properties. This process involves densification, pore elimination, and microstructural coarsening, primarily driven by diffusion and capillary forces. Once the liquid phase forms, it infiltrates grain boundaries, facilitating particle rearrangement, enhancing densification, and causing localized softening [181]. The resulting microstructure features bonded W grains embedded in a solidified matrix. Key parameters such as grain size, W volume fraction, contiguity, and dihedral angle govern the alloy’s performance [182].

In the early stages of sintering, densification occurs through bulk transport mechanisms, including grain boundary diffusion and solid-state activated sintering [8]. Grain boundary diffusion is the dominant densification mechanism between 1300 and 1750 °C. Upon the formation of the liquid phase, pre-existing W-W sinter bonds dissolve due to a solubility shift, enabling liquid penetration into grain boundaries. As sintering progresses, coarsening occurs through a dissolution–reprecipitation mechanism, where smaller grains dissolve into the liquid and reprecipitate onto larger ones, reducing the grain count and increasing the average grain size [183]. This grain growth is described by the cubic model:

$$G_{50}^3=G_0^3+Kt$$

(1)

where G₀ is the initial grain size, G₅₀ is the median grain size at the time t, and K is the temperature-dependent growth rate constant [184]. Pore elimination proceeds sequentially, with smaller pores shrinking first, promoting full densification [8]. In an 88W-10Ni-2Cu alloy sintered at 1412 °C, the median pore size initially increases due to grain restructuring and then decreases as full density is achieved [185]. The relationship between grain size and porosity is given by:

$$G_{50}=\theta G_0{(1-f)}^{-0.5}$$

(2)

where f is fractional porosity [57]. As sintering continues, grains adopt polygonal shapes, aiding in pore elimination and enhancing mechanical strength [186]. Although the presence of a liquid phase accelerates diffusion, excessive liquid content can hinder grain growth by increasing diffusion distance [187]. The grain growth rate parameter K (μm³/s) depends on the liquid volume fraction (V_L), as:

$$K=-1.1+1.9{V_L}^{2/3}$$

(3)

Experimental results reveal that higher liquid fractions reduce grain boundary interactions, moderating coarsening rates [188]. Additionally, liquid migration causes grain size variations along the sample height as heavier W grains settle and the liquid migrates upward. The transition from solid-state to liquid-assisted transport is critical to achieving optimal densification and mechanical integrity. Understanding these mechanisms is key to optimizing processing parameters for superior material performance.

The mechanical properties of WHAs are largely governed by microstructural features such as W grain size, contiguity, matrix volume fraction, and W solubility in the matrix [13]. The distribution, size, and morphology of these grains are influenced by sintering temperature, duration, and alloy composition. Grain size is commonly defined by the number of grains per unit area, and the yield strength is related to it via the Hall–Petch relationship [189].

$${\mathsf{\sigma }}_{\mathrm{y}}={\mathsf{\sigma }}_0+\mathrm{ }\mathrm{K}{\mathrm{d}}^{-1/2}$$

(4)

where σ_y is the yield strength, σ₀ the friction stress, K is the material constant, and d is the grain size. For WHAs, yield strength is also influenced by matrix volume fraction V_M and W grain size D, as shown by [99]:

$${\mathsf{\sigma }}_{\mathrm{y}}={\mathsf{\sigma }}_0+\mathrm{ }\mathrm{G}\mathrm{b}{\left({\displaystyle \frac{1-{\mathrm{V}}_{\mathrm{M}}}{\mathrm{D}{\mathrm{V}}_{\mathrm{M}}}}\right)}^{0.5}$$

(5)

where G is the shear modulus and b is the Burgers vector. Li et al. [65] proposed another model considering the W-W contiguity C_WW and matrix fraction, assuming that yielding begins in the Ni-Fe-W matrix:

$${\mathsf{\sigma }}_{\mathrm{y}}={\mathsf{\sigma }}_0+\mathrm{K}{\left({\displaystyle \frac{C_{\mathrm{W}\mathrm{W}}\left(1-{\mathrm{V}}_{\mathrm{M}}\right)}{{\mathrm{G}}_{\mathrm{W}}{\mathrm{V}}_{\mathrm{M}}}}\right)}^{0.5}$$

(6)

where G_W is the W grain size. This model suggests that yield strength increases with decreasing W grain size and matrix fraction. Ductility, according to Panchal and Nandy [30], depends on the matrix phase fraction and contiguity. A higher matrix fraction generally increases elongation due to the softer FCC structure of the matrix. W grain size also influences the failure mode during tensile testing. Large grains tend to fail via intergranular fracture—a lower energy process—whereas fine grains promote transgranular fracture due to the higher energy required to form new surfaces along W/matrix interfaces [14]. Humail et al. [66] found that increased W content reduces ductility due to higher W-W contiguity and fewer ductile matrix regions. The brittle W-W interfaces are more prone to microcrack formation and propagation, lowering the overall toughness. Contiguity (C_WW), which describes the connectivity between W grains, is calculated as:

$$C_{\mathrm{W}\mathrm{W}}={\displaystyle \frac{{2\mathrm{N}}_{\mathrm{W}\mathrm{W}}}{{\mathrm{ }\mathrm{ }2\mathrm{N}}_{\mathrm{W}\mathrm{W}}+{\mathrm{N}}_{\mathrm{W}\mathrm{M}}}}$$

(7)

where N_WW and N_WM are the counts of W-W and W/matrix interfaces, respectively, typically measured via line-intercept methods on SEM images [65]. Another critical microstructural parameter is the dihedral angle, the angle at the junction of two adjacent W grains. It impacts densification behavior: Low liquid and low dihedral lead to intermediate densification with distortion, while low liquid and high dihedral results in slow densification and minimal distortion. High liquid and low dihedral enable rapid densification with distortion, whereas high liquid and high dihedral causes slow densification with minimal distortion [8]. Optimizing liquid content and controlling the dihedral angle is essential for balancing densification, grain growth, and mechanical properties. Typically, the liquid phase content in LPS lies within the range of 5 to 15 vol.% [107].

The contiguity plots based on a comprehensive dataset compiled from 200 sources in the literature spanning a range of W contents underscore the critical influence of microstructural features on the mechanical behavior of WHAs fabricated through various sintering techniques (Figure 6a). The contiguity vs. W content plot reveals a consistent trend of increasing contiguity with increasing W content, reflecting the formation of a more interconnected W particle network. This effect is especially pronounced in conventionally sintered samples, where extended sintering durations and slower diffusion kinetics facilitate grain coarsening and promote extensive W-W contact. In contrast, SPS and microwave sintering result in lower contiguity at comparable W levels, primarily due to their rapid densification kinetics and the preservation of finer, more discrete W grains. Notably, SPS-processed samples exhibit slightly higher contiguity than the microwave-sintered counterparts at a similar W content. This distinction is attributed to the lower sintering temperatures typically employed in SPS, which limit diffusion and reduce W dissolution, thereby maintaining a greater degree of direct W-W contact. Post-sintering deformation processes, such as swaging or extrusion followed by heat treatment, further impact contiguity. These high-temperature mechanical and thermal treatments promote the dissolution of W into the matrix, disrupting the continuity of the W network and advancing the development of a more homogeneous microstructure. This reduction in contiguity enhances ductility and toughness by mitigating the brittleness associated with highly interconnected W frameworks while still preserving strength through improved matrix bonding and efficient load transfer.

Similarly, the W grain size vs. W content plot (Figure 6b) highlights the role of the sintering technique in controlling grain coarsening. Conventionally sintered WHAs show a clear increase in grain size with rising W content, primarily due to the extended thermal exposure that enables grain boundary migration and particle coalescence. In contrast, SPS and microwave sintering effectively suppress grain growth across the composition range due to the rapid densification and lower processing temperatures, resulting in finer and more uniform microstructures. Importantly, the data indicate that post-sintering deformation and subsequent heat treatments have minimal influence on W grain size, regardless of the initial sintering technique. This observation suggests that the inherent hardness and thermal stability of W grains limit further coarsening during mechanical processing and thermal exposure.

In addition, the dihedral angle vs. relative density plot (Figure 6c) reveals an inverse relationship, whereby the dihedral angle decreases with increasing relative density. This trend reflects enhanced densification during sintering, as higher relative densities are typically associated with the more effective liquid-phase wetting and capillary-driven rearrangement of W particles. Improved wetting behavior reduces the dihedral angle by strengthening solid–liquid interfacial interactions and minimizing pore volume at grain boundaries. As a result, lower dihedral angles at higher densities contribute to stronger interparticle bonding and a more cohesive microstructure, ultimately improving the mechanical performance.

Overall, these observations from the literature data highlight the role of sintering kinetics and densification mechanisms in controlling the key microstructural parameters, such as contiguity, grain size, and dihedral angle, that govern the mechanical response of WHAs.

5. Mechanical Properties of WHAs

WHAs’ mechanical properties depend on their W content (typically 80–97 wt.%), matrix composition, sintering parameters, and post-processing treatments. Increasing the W content enhances density and stiffness but can lead to grain coarsening and the formation of continuous W-W networks, which reduce ductility and promote brittleness. Achieving full densification is critical, as residual porosity significantly lowers both strength and ductility. In fully densified WHAs, hardness remains around 300 HV across a wide range of W contents but can reach ~428 HV with ceramic reinforcements (e.g., Y₂O₃, B₄C) or post-sintering treatments [8]. However, over-sintering causes grain growth, which degrades mechanical performance. In W-Ni-Fe systems with a 7:3 Ni:Fe ratio, the post-sintered matrix typically consists of ~54Ni-23Fe-23W, achieving a ~500 MPa tensile strength and ~133 GPa elastic modulus. The testing method affects the reported strength values; tensile testing offers the most reliable data while bending and compression often overestimate performance [190,191,192]. Optimal strength and ductility are typically achieved near 93 wt.% W, particularly after deformation-aging treatments (e.g., swaging + aging), which can raise hardness (~330 HV), yield strength (~791 MPa), and tensile strength (~1039 MPa). Above 95 wt.% W, grain coarsening and increased contiguity degrade performance, and tensile strength becomes less consistent due to microstructural variability [193].

Processing techniques significantly influence WHA performance. Processing strongly influences mechanical behavior. Annealing can reduce strength by 15–20% but typically increases ductility by 20–30%. Deformation-aging (10–25% swaging + heat treatment) can increase tensile strength to 1375–1435 MPa, with an elongation of 7–10% [8]. Optimized sintering and deformation-aging are essential for balancing strength and ductility. Further, WHAs’ performance is governed by phase morphology, volume fraction, and interfacial characteristics. Weak W/W and W/matrix interfaces often cause premature failures in the matrix, while strong interfacial bonding shifts failure to W cleavage, improving strength. The first WHAs were based on W-Ni-Cu, and while they demonstrated higher density, they suffered from lower strength and corrosion resistance than W-Ni-Fe systems. W-Ni-Fe remains the dominant system, with 7:3 or 8:2 Ni–Fe ratios offering good mechanical balance. Excessive alloying may promote intermetallic phases and embrittlement [14]. Wu et al. [24] compared the LPS behavior of 80W-14Ni-6Cu vs. 80W-14Ni-6Fe alloys. While both reached full densification, the Ni-Cu alloy distorted significantly at higher sintering temperatures, whereas the Ni-Fe alloy retained dimensional stability up to 1500 °C. Kiran et al. [99] showed that adding 1.6 wt.% Co to 93W-4.9Ni-2.1Fe increased W solubility in the matrix and enhanced strength. Similarly, 0.2 wt.% Re addition further improved tensile strength.

Typically, sintered W-Ni-Cu/W-Ni-Fe often exhibit a strength range of 0.7–1 GPa, below the >1 GPa required for KEPs [70]. Das et al. [58] showed that a 91W-7Ni-1.5Fe-0.5Co alloy had superior room-temperature mechanical properties due to finer grains and lower porosity. Post-sintering heat treatment improved tensile behavior, and a 5Ni-5Fe matrix enabled >16% elongation without compromising strength. Kumari et al. [79] found that increasing the Ni/Co ratio in W-Ni-Co alloys suppressed fine W precipitates, softening the alloy. Swaged WHAs showed increased hardness with lower Ni–Co ratios. Panchal et al. [68] studied the flow and work-hardening behavior of 92W-5.5Ni-2.5Fe, noting that swaged specimens had higher strength, while heat-treated ones showed better elongation. Humail et al. [66] observed that 93 wt.% W alloys achieved the highest tensile strength due to optimized grain size, though ductility declined at higher W contents.

W-Ni-Co is a relatively newer ternary alloy system than W-Ni-Fe and W-Ni-Cu alloys. Co is commonly added to improve strength and ductility by reinforcing the binder and enhancing W grain wettability and interfacial bonding. Dincer et al. [194] investigated the microstructures of the W-Ni-Co system and found that microstructures consisted of rounded, almost pure W grains and a Ni-Co-W binder matrix phase. Adding Co led to an increase in W solubility in the matrix phase of up to 42 wt.%. However, higher Co content promoted the formation of brittle intermetallic phases (e.g., Ni₂Co₃W, Co₇W₆), reducing toughness [80]. Skoczylas et al. [82] found that a 91W-6Ni-3Co alloy reached theoretical density (~17.46 g/cm³), and that swaging from 15% to 20% increased the strength by ~70 MPa. Despite its mechanical benefits, Co poses health risks such as lung fibrosis and asthma [195], prompting efforts to develop Co-free WHAs. Cury et al. [196] demonstrated that Co-free alloys, with optimized Ni/Fe ratios and heat treatment, can match the mechanical performance of Co-containing variants. In ballistic testing, Co-free alloys achieved comparable penetration, confirming their suitability for KEPs. Further improvements have been made by refining W-Ni-Fe microstructures through rare earth additions (Y, La, oxides) and refractory metals (Re, Ta, Mo) [197]. Powder mixing methods also influence the mechanical properties. Eroglu et al. [94] compared wet attritor milling and dry turbula mixing for 90W-7Ni-3Fe and 92.5W-5.25Ni-2.25Fe alloys, finding that wet-milled powders produced denser, more ductile alloys for a given W powder size. Coarser W particles increased ductility in turbula-mixed alloys, whereas finer powders resulted in the lowest tensile strengths.

Figure 7 collectively illustrates the influence of different sintering techniques—conventional, microwave, and SPS—on the mechanical properties of WHAs, based on an extensive dataset compiled from over 250 sources in the literature, spanning a wide range of W contents. Both microwave and SPS methods generally produce higher tensile strengths compared to conventional sintering. This improvement is primarily attributed to their rapid heating rates and enhanced densification kinetics, which facilitate the formation of finer and more homogeneous microstructures. In contrast, WHAs processed via conventional sintering typically exhibit lower as-sintered strengths. However, post-sintering treatments such as swaging or extrusion followed by heat treatment play a critical role in improving tensile properties by refining the microstructure and reducing residual porosity, thereby enabling conventionally sintered alloys to reach performance levels comparable to those achieved through advanced sintering routes. Elongation data show greater variability, particularly in SPS-processed samples. This variability is likely due to the lower sintering temperatures associated with SPS, which make ductility more sensitive to microstructural fluctuations and W content. Conventionally sintered alloys generally exhibit lower elongation, often due to coarser W grains and higher porosity, unless improved through secondary processing. Similarly, both SPS and microwave sintering typically yield higher hardness values due to the improved densification and refined grain structures. Despite these general trends, all three datasets show substantial standard deviations, highlighting the complex interplay between processing conditions, alloy composition, and post-sintering treatments.

The contiguity plots further emphasize the role of microstructural connectivity in influencing mechanical behavior across different sintering techniques (Figure 7d,e). In the contiguity versus elongation plot, a clear negative correlation is observed: higher contiguity tends to reduce elongation, as increased W-W contact promotes early strain localization and brittle fracture. The contiguity vs. tensile strength relationship shows more complex behavior. While moderate contiguity contributes to strength through better load transfer across the W network, excessively high contiguity (common in high W alloys) leads to brittleness and premature failure. Additionally, the graph depicting the dihedral angle versus tensile strength reveals an inverse relationship: as the dihedral angle decreases, tensile strength increases (Figure 7f). This trend is primarily linked to the wetting behavior of the liquid binder phase (typically Ni-Fe or Ni-Co) during LPS. A lower dihedral angle reflects the improved wetting of W particles by the matrix, which strengthens the interfacial bonding between the hard W grains and the ductile binder. Stronger interfaces enable more effective load transfer and reduce the risk of intergranular fracture. Moreover, improved wetting promotes a more uniform and defect-free microstructure, further contributing to enhanced tensile strength. Conversely, a higher dihedral angle indicates poor wetting, resulting in weak interfacial bonding, increased porosity, and a higher susceptibility to crack initiation and propagation. Thus, optimizing sintering conditions to achieve lower dihedral angles is essential for improving the mechanical integrity of WHAs.

Summary of mechanical properties of WHAs with different sintering methods and with different alloying elements are presented in

Table 8. Summary of mechanical properties of WHAs with different sintering methods.

Composition (wt.%)	Sintering Method	Contiguity	W Particle Size	UTS (MPa)	Elong. (%)	Remarks	Ref.
92.5W-6.4Ni-1.1Fe	Conventional sintering at 1500 °C	0.32	17.3	642	3.5	Microwave-sintered specimens give fine W grains, resulting in higher strength.	[61]
92.5W-6.4Ni-1.1Fe	Microwave sintering at 1500 °C	0.42	9.4	805	11.2
90W-7Ni-3Fe	Conventional sintering at 1480 °C, heating rate of 5 °C/min	-	-	862	19.7		[28]
90W-7Ni-3Fe	Microwave sintering at 1480 °C heating rate of 80 °C/min	-	-	922	19.7
90W-7Ni-3Fe	Microwave sintering at 1460 °C	0.26	21	740	9		[27]
90W–7Ni-3Fe	Conventional sintering at 1460 °C	0.51	32	650	5
90W-6Ni-2Fe-2Co	Microwave sintering at 1470 °C	0.30	21	989	16	Microwave-sintered specimens give lower contiguity and higher ductility.
90W-6Ni-2Fe-2Co	Conventional sintering at 1460 °C	0.62	47	682	4
90W-7Ni-3Fe	SSS at 1350 °C	-	2.2	1204	-	SSS gives higher tensile strength because of fine W grains.	[26]
90W-7Ni-3Fe	SPS at 1100 °C, heating rate of 100 °C/min, pressure of 50 MPa	-	-	930	2.91	SPS gives higher tensile strength but lower ductility.	[67]

and

Table 9. Summary of mechanical properties of WHAs with different alloying elements.

Composition (wt.%)	Sintering method	Contiguity	W Particle Size	UTS (MPa)	Elong. (%)	Remarks	Ref.
92W-5.6Ni-2.4Fe	SPS, sintered at 1400 °C	0.64	12.3	975	12	Increasing the Co content increases the strength. A total of 1 wt.% Co gives the highest tensile strength.	[13]
91.5W-5.6Ni-2.4Fe-0.5Co	SPS, sintered at 1400 °C	0.54	11.56	961	16
91W-5.6Ni-2.4Fe-1.0Co	SPS, sintered at 1400 °C	0.47	9.48	1508	20
90.5W-5.6Ni-2.4Fe-1.5Co	SPS, sintered at 1400 °C	0.43	9.68	1330	21
90W-5.6Ni-2.4Fe-2.0Co	SPS, sintered at 1400 °C	0.50	11.1	1256	18
93W-4.9Ni-2.1Fe	Conventional sintering at 1480 °C	0.53	56	858	17	Adding Co gives a slightly higher strength. Re refines the W grains.	[99]
93W-4.2Ni-1.2Fe-1.6Co	Conventional sintering at 1480 °C	0.43	54	869	18
93W-4.9Ni-1.9Fe-0.2Re	Conventional sintering at 1480 °C	0.48	49	876	18
93.5W-4Ni-2Fe-0.5Co	Conventional sintering at 1500 °C	-	-	853	14	Adding Mo improves ductility due to a larger volume fraction of matrix.	[14]
90.5W-7.2Ni-1.8Fe-0.45Co-0.05Mo	Conventional sintering at 1500 °C	-	-	950	21
95W-3.5Ni-1.5Cu	Conventional sintering at 1510 °C	0.6	60	660	3	Tensile properties and hardness of the W-Ni-Cu alloy are inferior due to the coarser W grains.	[58]
96W–3Ni–1Cu	Conventional sintering at 1510 °C	0.7	70	660	3

, respectively. Overall, while microwave and SPS techniques provide superior strength and hardness, conventionally sintered WHAs, when subjected to appropriate post-sintering treatments, can achieve comparable tensile and ductility characteristics. The significant data scatter highlights the need for careful optimization of processing routes to reliably tailor the mechanical performance of WHAs for critical applications.

6. Deformation Behavior of WHAs

The deformation behavior of WHAs is complex and governed by the interplay between W grains and the matrix phase. Key factors such as temperature, strain rate, and microstructural characteristics significantly affect their mechanical response. Understanding these aspects is essential for optimizing WHAs for various applications. Kim et al. [198] examined how the W particle shape affects the dynamic deformation and fracture behavior in WHAs using Kolsky bar torsional tests. Microstructural characterization (e.g., micro-CT, fractography), mechanical property measurements, and fractography results revealed that WHAs produced with coarse and irregularly shaped W particles contributed to a predominantly cleavage fracture mode. In contrast, WHAs with finer particles showed more localized central fractures with minor shear deformation, indicating different deformation and failure mechanisms that aligned closely with their observed impact behavior.

Park et al. [177] explored dynamic deformation in ODS WHAs, incorporating 0.1 wt.% Y₂O₃ via mechanical alloying. The mechanical alloying process refines the W particle size without increasing the interfacial area fraction, promoting interfacial de-bonding and ASB formation. The ODS WHAs exhibited lower maximum shear stress and fracture strain than conventionally sintered specimens, suggesting improved penetration performance through enhanced self-sharpening. Fang et al. [32] investigated the penetration and self-sharpening behavior of 80W-14Cu-6Zn, 90W-7Ni-3Fe, and 35CrMnSi alloys. Their results showed that the 80W-14Cu-6Zn penetrator benefited from high strength and an appropriate critical failure strain. They proposed that the melting and ejection of the Cu-Zn matrix during penetration acts as a lubricant, reducing lateral force and enhancing self-sharpening. Ryu et al. [199] studied 93W-5.6Ni-1.4Fe WHAs fabricated by mechanical alloying and sintering (1445–1485 °C, H₂ atmosphere). They found that increasing the matrix content and reducing W/W contiguity enhanced the impact energy. The processing method also significantly affected the failure mode, which shifted from a brittle intergranular fracture in SSS WHAs to a ductile matrix shear fracture in LPS WHAs. These findings emphasize the importance of promoting a high matrix volume and low W contiguity for optimal mechanical performance [75].

Gong et al. [200] analyzed the dynamic deformation of a fine-grained 93W-4.9Ni-2.1Fe-0.03Y (wt.%) alloy with an average grain size of 6 µm at different strain rates. At higher strain rates, ASBs containing dense dislocations and twins were observed despite minimal thermal softening, suggesting that adiabatic shear failure can occur independently of a significant temperature rise. Grain refinement and Y-containing second-phase particles altered the deformation mechanisms of fine-grained WHAs relative to coarse-grained WHAs.

Self-sharpening in WHAs is critical for improving penetration performance and relies on localizing ASBs under high shear strain. U alloys, especially U-3/4Ti, historically excel due to ASB-induced self-sharpening, but environmental risks have shifted attention toward alternatives like WHAs and U-6Nb alloys. Despite WHAs’ high melting point and strain-rate sensitivity, they often suffer from mushrooming upon impact, which reduces penetration depth [3,201]. Mushrooming refers to the deformation mode where the leading end of a projectile flattens and expands radially upon high-velocity impact, resembling the cap of a mushroom. This phenomenon is commonly observed in penetrators made of pure W and WHAs, particularly under conditions where adiabatic shear localization is delayed or suppressed. W has a very high melting point (~3420 °C) and low thermal softening, meaning that even under intense plastic deformation, it resists flow softening. Moreover, W exhibits strong strain-rate sensitivity, so as the strain rate increases, the material’s resistance to deformation increases as well. These characteristics delay the onset of localized plastic flow (i.e., shear banding), and instead, the material deforms more uniformly and radially at the impact face. This deformation mode reduces the penetration efficiency, as energy is dissipated in a radial expansion rather than being concentrated along a narrow axial path. In contrast, uranium alloys, such as U-3/4Ti, exhibit early shear localization and fracture, producing a self-sharpening chiseled nose that maintains a narrow cross-section and penetrates deeper into the target [6].

Key strategies to mitigate this issue include reducing the W/W interfacial area and increasing the W grain size to lower the contiguity. Various testing methods, such as Hopkinson bar tests [202], normal plate impact tests [203,204,205], and oblique flyer-plate impact experiments [206], have been employed to study WHA behavior under dynamic loading. ODS-WHAs (0.1 wt.% Y₂O₃) demonstrate enhanced ASB formation due to improved interfacial bonding [177], while pre-twisted WHAs exhibit W grain aspect ratio and orientation-dependent ASB formation [207]. ASBs nucleate where maximum shear stress aligns with the major axis of W grains and propagate through the Ni-Fe matrix, which requires lower shear deformation energy.

Microstructural heterogeneity, particle shape, and matrix content significantly influence ASB formation [198,208]. Ultrafine-grained (UFG) W (~500 nm) exhibits lower strain-rate sensitivity and pronounced flow softening at ~10³ s⁻¹ strain rates, with peak stresses around 3 GPa. The strain-rate sensitivity of UFG W has been reported to be lower than that of coarse W [209]. WHA penetrators processed via hydrostatic extrusion and hot torsion (HE + HT) exhibit enhanced self-sharpening behavior, maintaining an acute shape upon ballistic impact due to localized ASBs, unlike as-sintered or as-extruded WHAs, which develop mushroomed heads [210]. Temperature and strain rate also play vital roles. The deformation behavior of W-Ni-Fe alloys has been studied across a range of temperatures (25–1100 °C) and strain rates (8 × 10²–4 × 10³ s⁻¹) [211]. W-Ni-Fe WHAs shows increasing flow stress with strain rate and decreasing work hardening with temperature. At high shearing strain rates (~7 × 10⁵ s⁻¹), pressure-shear plate impact tests show thermal-softening-induced ASBs, demonstrating significant strain-rate sensitivity [212,213]. Using Taylor impact tests, ASB formation has been quantified as a function of impact velocity [214]. Reverse ballistic impact tests at velocities of 173 m/s and 228 m/s reveal a slight ellipticity in the mushroomed end, with low-speed impacts producing radial cracks and higher speed impacts leading to localized deformation along a curved path [215]. Numerical simulations using Johnson–Cook constitutive models provide equivalent failure strains for ASB formation [216].

Shear banding in WHAs often originates at the transition between the mushroomed region (deformed) and the undeformed regions. ASBs occur across various temperatures. At cryogenic temperatures, less shear strain occurs, indicating that temperature influences ASB formation. Shear banding has been observed in liquid-phase-sintered 90W-7Ni-3Fe alloys at strain rates of 5400 s⁻¹ [217]. High-strain-rate Kolsky bar tests confirm that WHA deformation is primarily governed by W grain behavior [202].

The thermo-mechanical mismatch between W particles and ductility in WHAs promotes conditions favorable for shear localization. W particles exhibit higher flow stress, density, strain-rate sensitivity, and thermal conductivity, along with lower strain hardening and specific heat, compared to the matrix. As a result, the matrix deforms earlier under load, generating heat that is rapidly conducted into the W particles. This raises their temperature, reduces their flow stress, and promotes ASBs. This thermo-mechanical coupling explains why composite WHAs are more susceptible to shear localization than their individual constituents [218]. ASBs are narrow zones of intense plastic deformation caused by localized adiabatic heating at high strain rates, such as those experienced during ballistic impact. Under such conditions, the rate of heat generation exceeds the rate of thermal dissipation, triggering a rapid temperature rise and localized thermal softening. If the deformation fails to be controlled, it leads to a “mushrooming effect”, which increases the diameter of the penetrator’s head and reduces its ability to penetrate [3].

Gong et al. [200] observed <10 µm wide ASBs in fine-grained WHAs under a 1.9 × 10³ s⁻¹ strain rate. Wu et al. [219] demonstrated ASB formation in 93W-4.5Ni-1.5Fe-1Co under high-temperature split Hopkinson pressure bar and hypervelocity impacts into a concrete target test using a two-stage light gas gun. They demonstrated ASBs as fracture initiation sites and a key failure mode in WHAs under dynamic loading but also showed that they contribute to self-sharpening, a desirable behavior that enhances penetration efficiency in armor-piercing applications [220].

W grain size and matrix composition strongly influence ASB formation and penetration performance [200]. Luo et al. [221] showed that fine-grained 95W rods achieved greater penetration depth and produced narrower craters compared to coarse-grained rods, which exhibited mushrooming and wider craters. Kim et al. [63] revealed that ASB morphology was affected by W particle elongation, interfacial strength, and matrix composition. Cracks often initiated in shear bands but sometimes diverted along W/W interfaces due to their brittleness, which hindered shear band growth. Therefore, reducing W/W particle interfaces through microstructural control and optimized fabrication (e.g., heat treatment) was essential to promoting narrow, high-strain ASBs that facilitated crack propagation and self-sharpening.

An effective strategy to improve the ASB behavior in WHAs is to use alternative matrix materials with high susceptibility to shear localization. BMGs [86], HEAs [88,222], and systems like W-Ni-Mn [173] and W-Cu-Zn [223] have shown potential. Ma et al. [86] fabricated WHA-BMG composites using SPS with varying W volume fractions in a Zr₅₅Cu₃₀Al₁₀Ni₅ matrix. Under compressive loading, elastic modulus mismatch at the W/matrix interface generated stress concentrations that initiated shear bands. A higher W content promoted band formation, increasing free volume and facilitating localized plasticity. Hu et al. [87] demonstrated that W-reinforced, Zr-based BMGs (Zr₅₅Cu₃₀Al₁₀Ni₅) retained a sharp tip during penetration due to matrix softening and W plastic deformation, effectively preventing mushrooming.

Despite their advantages, BMGs suffer from limited macroscopic plasticity and poor fracture toughness [86]. Processing BMGs is also a challenge due to the critical cooling rate requirements to retain the amorphous phase; methods like spark plasma sintering are often required, limiting their scalability. Although high-density BMGs based on elements like Hf and Ta offer improved ductility and density, economic and processing constraints hinder their broader application in penetrators [224]. Conventional WHAs like W-Ni-Mn [173] and W-Cu-Zn [223] systems also enhance shear localization. Mn reduces the thermal conductivity of the WHAs and leads to enhance shear localization, while a Cu-Zn matrix facilitates shear failure. However, both systems suffer from drawbacks: Mn high vapor pressure increases porosity, and Cu-Zn soft matrix reduces dynamic strength, thereby necessitating the development of new matrix materials that overcome these limitations [225].

Recently, HEA as a matrix in WHAs has shown promise in improving penetration performance [210]. HEAs represent a novel category of materials that consist of several metallic elements, usually four or more, combined in approximately equal atomic ratios [226,227]. Liu et al. [228] developed a W-FeNiMo HEA-WHA featuring BCC, an FCC matrix, and μ-phase precipitates. These precipitates promoted dynamic recrystallization via localized strain gradients, sustaining sharp projectile heads and improving penetration depth by 10–20%. Zhou et al. [88] demonstrated that an AlCrFeNiV HEA matrix improved W/matrix interface properties. HEAs with low W solubility and diffusivity were particularly effective in refining W grains. Satyanarayana et al. [89] used CoCrFeMnNi HEAs as a crack-resistant WHA matrix. In addition to W grains and the HEA phase, consolidated WHA composites also showed the existence of a Cr-Mn-rich oxide phase. The Cr-Mn-rich oxide phase restricted the W grain growth. The SPS-processed composites exhibited higher compressive strength (~2041 MPa). Panigrahi et al. [90] reported improved densification, hardness, and strength in 90W-10FeNiCoCrCu HEA-WHAs. Anwar et al. [222] further showed that HEA-WHAs delivered 42% higher hardness and enhanced strain hardening, supporting superior penetration performance and positioning them as a viable alternative.

7. Modeling and Simulation of WHAs

Modeling and simulating WHAs requires a multidisciplinary approach that combines materials science, mechanical engineering, and computational techniques. Given their application in high-strain-rate environments, understanding the dynamic mechanical behavior of WHAs is crucial for optimizing performance. Sharma et al. [229] developed four constitutive models, the Modified Johnson–Cook (m-JC), Modified Zerilli–Armstrong (m-ZA), Modified Arrhenius (m-Arr), and Modified Khan–Huang–Liang (m-KHL) models, to predict the flow stress behavior of a 92.5W-5.25Ni-2.25Fe alloy. These models accounted for the coupled effects of strain, strain rate, and temperature. Processing maps were constructed at strains of 0.16, 0.18, and 0.20. Optimal processing conditions were identified at 0.16 strain, high strain rates (2500–4000 s⁻¹), and temperatures between 427 and 627 °C, making this ideal for evaluating WHA penetration performance.

Finite element modeling (FEM) plays a pivotal role in simulating sintered component behavior, accounting for surface tension, gravity, grain growth, friction, and solid content. Park et al. [230] used FEM with constitutive equations for linear viscous compressive materials to simulate grain evolution, densification, and distortion. Li et al. [231] conducted detailed FEM simulations to investigate the penetration performance of W fiber/Zr-based metallic glass composite rods, comparing them to conventional WHA rods. The simulations incorporated a microstructure-based geometric model and a modified coupled thermo-mechanical constitutive model to accurately reflect the high strength and shear sensitivity of the metallic glass matrix under high-strain-rate conditions. Li et al. [232] created W/BMG composites using LS-DYNA, incorporating realistic particle distributions via a Monte Carlo approach. Couque et al. [214] examined how shear banding affects the penetration performance of four conventional WHAs, each containing 90 wt.% W, a ~17.1 g/cm³ density, and a ~1100 MPa yield stress, but differing in ductility and impact energy. Using symmetric Taylor impact tests (200–310 m/s) and ballistics tests (~1400 m/s against 303–330 HB steel) and AUTODYN-2D simulations with J-C models, they identified the critical impact velocities (199.8–223.2 m/s) for shear band initiation and the corresponding equivalent plastic strains.

Kennedy and Murr [233] provided a benchmark for simulation accuracy by comparing measured residual microhardness maps and penetration channel geometries with simulations generated via AUTODYN-2D. Hafizoglu and Durlu [234] highlighted that smoothed-particle hydrodynamics (SPH)-based simulations more accurately predicted the complex fracture shapes and dimensions observed in Taylor impact tests than mesh-dependent FEM approaches. FEM and SPH have proven essential for simulating high-strain-rate deformation and failure modes, including fragmentation and adiabatic shear banding. The Johnson–Cook model remains effective for WHAs, while modified coupled thermo-mechanical models address the shear localization and amorphous nature of metallic glasses [232,234]. Hafizoglu et al. [235] demonstrated that increasing the sintering temperature to 1480 °C and adjusting the Ni/Fe ratio to 4.67/2.33 optimized the plasticity and decreased contiguity, leading to improved fragmentation and energy transfer during impact. Simulation results correlated well with experimental data in terms of deformation at the impact and rear surfaces of targets, indicating that material design can be finely tuned to maximize ballistic effectiveness.

Lee et al. [211] examined the effect of strain rate and temperature on the impact deformation of a 92.5W-5.25Ni-2.25Fe LPS WHA. They employed the Zerilli–Armstrong dislocation mechanics-based model to describe the material’s rate-dependent behavior, which successfully captured the deformation response under high-rate and high-temperature conditions. This model is particularly valuable as it distinguishes between body-centered cubic (BCC) and face-centered cubic (FCC) metals. The mechanical performance of WHAs depends on the characteristics of each phase and the strength of interfacial bonding. Stress redistribution is limited under high-strain-rate loading, such as in explosions, leading to differences in stress and strain across phases. This mismatch contributes to damage mechanisms, including W grain cleavage, the ductile fracture of the matrix, and interfacial separations (W/W and W/matrix) [236]. Tang et al. [237] developed a two-phase constitutive model with damage for a 92.5W-4.9Ni-2.1Fe-0.5Co alloy. Using the Mori–Tanaka mean field theory and Eshelby’s equivalent inclusion theory, they incorporated damage evolution through continuum damage mechanics. The model showed good agreement with experimental data, except at extreme strain rates.

WHAs are modeled as two-phase composites consisting of rigid W particles dispersed within a ductile Fe-Ni-W matrix. Each phase is described using the Johnson–Cook (J-C) constitutive model, incorporating thermal softening based on the Zhou et al. formulation [238]. For the W particles, the material parameters are as follows: yield stress A = 730 MPa, hardening modulus B = 562 MPa, strain-rate sensitivity C = 0.029, strain-hardening exponent n = 0.0751, shear modulus μ = 155 GPa, bulk modulus K = 317 GPa, density ρ = 19,300 kg/m³, and specific heat capacity c = 138 J/kg·°C. The Fe-Ni-W matrix is characterized by the following: A = 150 MPa, B = 546 MPa, C = 0.0838, n = 0.208, μ = 98.84 GPa, K = 202.4 GPa, ρ = 9200 kg/m³, and C = 382 J/kg·°C. Taylor impact simulations conducted on a cylindrical rod measuring 6.35 mm in diameter and 25.4 mm in length at impact velocities between 173 and 228 m/s produce deformation results that closely align with experimental observations [238].

Sun et al. [239] evaluated a 90W-7Ni-3Fe WHA under both quasi-static (0.001–0.1 s⁻¹) and dynamic (2000–6000 s⁻¹) compression conditions, providing valuable experimental validation for simulation models. The constitutive behavior was accurately captured using the J-C model with the parameters A = 730 MPa, B = 1614 MPa, n = 0.16, and C = 0.109.

The J-C model remains widely used due to its simplicity and limited calibration requirements. It captures strain, strain rate, and temperature effects, and differentiates between BCC and FCC materials [240]

$${\sigma }_{\mathrm{y}}=(A+B{ϵ}^n)\left(1+Cln{\displaystyle \frac{{\mathsf{\epsilon }}^{\dot{}}}{{{\mathsf{\epsilon }}^{\dot{}}}_0}}\right)\left(1-{\left({\displaystyle \frac{T-T_r}{T_m-T_r}}\right)}^m\right)$$

(8)

where ε is equivalent plastic strain, ${\epsilon }^{\dot{}}$ is plastic strain rate, T is absolute temperature, $T_r$ is reference temperature, $T_m$ is melting temperatures, and A, B, n, C, and m are material constants.

The Zerilli–Armstrong (Z-A) model is based on dislocation mechanics and includes strain hardening, strain-rate sensitivity, and thermal softening. It also allows for discrimination between BCC and FCC materials and coupling between temperature and strain rate [240]:

$$\mathrm{ }{\sigma }_{\mathrm{y}}=C_{1\mathrm{ }}+C_{2\mathrm{ }}{e}^{\left(-C_{3\mathrm{ }}+C_{4\mathrm{ }}ln{\mathsf{\epsilon }}^{\dot{}}\right)\mathrm{T}}+C_{5\mathrm{ }}{ϵ}^n$$

(9)

Here, C₁ relates to dislocation mechanisms and grain size, and C₂-C₅ and n are material constants. Scapian [240] applied both J-C and Z-A models to a W-Ni-Cu alloy with the following parameters: ${{\epsilon }^{\dot{}}}_0=1{\mathrm{s}}^{-1}$, $T_r=298\mathrm{K}$, and $T_m=1673\mathrm{K}$. The optimized parameters were A = 647 MPa, B = 1337 MPa, n = 0.5616, C = 0.0640, and m = 0.4937. Both models showed excellent correlation with experimental flow stress data (R² = 0.9883). The dynamic materials model treats the workpiece as an energy dissipation system. Under steady strain and temperature, the input power (P) during plastic deformation is split into two components: G (heat from plastic deformation) and J (internal energy change due to microstructural evolution) [229]:

$$P=\sigma {\mathsf{\epsilon }}^{\dot{}}={\int }_0^{{\epsilon }^{\dot{}}}\sigma {\mathsf{\epsilon }}^{\dot{}}+{\int }_0^{\sigma }{\epsilon }^{\dot{}}d\sigma =G+J$$

(10)

The strain-rate sensitivity (m) is defined as:

$$m={\displaystyle \frac{\partial J}{\partial G}}={\displaystyle \frac{\partial \left(ln\sigma \right)}{\partial \left(ln{\epsilon }^{\dot{}}\right)}}$$

(11)

m influences the energy distribution between G and J. BCC metals exhibit higher m values due to the limited mobility of screw dislocations, leading to increased flow stress [239]. Jinzhu et al. [241] conducted experimental and numerical studies on WHA projectile impacts against alumina armor. Their simulations accurately predicted penetration depth and failure patterns, aligning well with experimental results.

Crystal plasticity-based simulations offer critical insight into the role of microstructure in the dynamic response of WHAs. These models incorporate grain-level anisotropy and account for variations in grain orientation and boundary strength, both of which significantly influence damage localization and spall resistance. By integrating cohesive zone models at grain and phase boundaries along with rate-dependent constitutive laws, the simulations successfully replicate experimentally observed, free-surface velocity profiles and spall fracture morphologies [242,243,244]. Experimental data show the spall strength of WHAs between 1.7 and 2.0 GPa, reflecting the influence of factors such as grain structure, binder distribution, and impact conditions [242]. Fracture initiates once microcracks grow to a critical size, typically through a combination of cleavage within W grains and separation along grain boundaries. During plate impact tests, a shock wave travels through the material and reflects off the free surface as a rarefaction wave, generating intense tensile stresses. These stresses are especially high at interfaces, where spall fracture is likely to begin. The slope of the stress–time trajectory, from the shocked to the released state, serves as an indicator of spall strength, which increases with higher impact velocities [244]. Microstructural features and loading direction play key roles in determining how damage evolves. Cracks tend to initiate along weaker paths, such as grain boundaries or soft matrix regions, before coalescing into larger fracture zones. Post-impact analysis consistently reveals both intergranular cracking and cleavage, confirming the dual fracture mechanisms in WHAs. Compared to DU, WHAs generally exhibit more complex damage patterns and higher resistance to spall failure [242,243].

Collectively, these studies demonstrate the critical role of validated constitutive modeling and multi-scale simulation in predicting WHA performance under extreme conditions. Models that incorporate microstructural detail, interfacial behavior, strain rate, and temperature sensitivity provide a robust framework for designing next-generation KEPs.

8. Conclusions and Future Work

WHAs have emerged as critical materials for KEPs, primarily due to their high density, strength, and ductility. This review has summarized recent developments in the processing, microstructural control, mechanical property enhancement, and deformation behavior of WHAs. Among the various consolidation techniques, LPS remains the most widely adopted due to its ability to achieve near-theoretical densification. Nevertheless, challenges such as W grain coarsening, elevated W-W contiguity, and heterogeneous microstructures continue to limit the mechanical performance of WHAs.

Considerable efforts have been directed toward optimizing powder preparation, sintering parameters, and alloy chemistry to address these limitations. Alloying additions such as Co, Mo, Re, and rare earth oxides (e.g., Y₂O₃, La₂O₃, CeO₂) have shown promise in enhancing matrix strength, refining microstructures, and promoting shear localization, an essential factor for improving the self-sharpening behavior of WHAs. Moreover, post-sintering treatments such as hot isostatic pressing, heat treatment, and rotary swaging have demonstrated significant improvements in tensile strength, hardness, and impact resistance by eliminating porosity and refining the grain structure.

Despite these advancements, WHAs still exhibit inferior penetration efficiency compared to DU alloys, primarily due to their mushrooming failure mode upon impact. While high-strain-rate ductility and strength have improved through microstructural engineering, the intrinsic limitations associated with W high melting point, brittleness, and interfacial cohesion remain critical challenges.

Future research should focus on advancing WHA systems through the following directions:

• Development of matrix compositions and alloying strategies that enhance thermal softening and shear localization without compromising ductility.
• Design and integration of HEA or multi-principal-element alloy matrices for improved bonding and thermal stability.
• Optimization of AM parameters to overcome defects such as porosity and cracking, enabling the fabrication of net-shape, high-performance WHA components.
• Use of ODS alloys to promote fine grain size, inhibit crack propagation, and enhance high-temperature performance.
• Implementation of data-driven approaches and modeling frameworks to predict phase evolution, densification behavior, and dynamic failure mechanisms under impact conditions.

By addressing these challenges, WHAs can be further engineered to approach or even surpass the terminal ballistic performance of DU alloys, thereby offering a safer, sustainable, and high-performance solution for next-generation penetrator applications.