Manufacturing of Diamond Tool Segments via Microwave–Hybrid Sintering

Abstract

Microwave (MW) sintering offers a promising alternative to conventional heating in powder metallurgy, providing faster processing, lower energy consumption, and improved microstructural control. In the diamond tool industry—where cost-efficiency and competitiveness are critical—MW–hybrid sintering shows strong potential for producing segments designed for cutting and polishing natural stone and construction materials. This study investigates the effects of sintering temperature, dwell time, and green density on the densification and mechanical properties of metal matrix composite (MMC) segments containing diamond particles. MW sintering reduced the optimum sintering temperature by 90–170 °C when compared to conventional free sintering. Under optimal conditions (57% green density, 820 °C, 5 min dwell), segments achieved ~95% densification and mechanical properties comparable to hot-pressed (HP) samples. Although MW–hybrid sintered matrices exhibited slightly lower Young’s modulus (~15%) and Vickers hardness (~20%), their flexural strength and fracture toughness remained comparable to HP counterparts. Overall, MW hybrid sintering provides a cost-effective, energy-efficient, and scalable route for fabricating high-performance diamond tool segments, supporting both economic viability and sustainable, competitive manufacturing.

1. Introduction

The diamond tool manufacturing industry is currently facing mounting pressure due to falling market prices and increasing global competition—particularly from low-wage regions—and tightening technical and environmental regulations [1]. These challenges underscore the urgent need for innovations that improve both the performance and sustainability of diamond tools. Current research efforts focus on reducing production costs, enhancing the wear resistance of metallic matrices, and optimizing diamond distribution within tool segments [2].

The development of synthetic diamonds in the 1950s catalyzed major advances in diamond tool technology. Today, diamond-impregnated tools (DITs) are widely used for stone cutting, drilling, machining, and geological exploration [3]. In ornamental stone processing, tools such as segments or beads consist of composite structures—metallic matrices embedded with 5–10 vol.% synthetic diamonds—typically welded onto a steel core [4]. These components are commonly fabricated via powder metallurgy (PM), involving the compaction of metal powders followed by hot pressing (HP) below the solidus temperature [5].

Circular saw blades, wire saws, and core drills represent the largest segment of DITs, particularly in natural stone processing and civil engineering applications [1]. For optimal cutting performance, the metallic matrix must retain the diamond particles effectively while wearing at a controlled rate to continually expose sharp edges. Matrix composition must therefore be carefully tailored to the tool type, operating conditions, and target materials [6]. While tool selection remains complex, higher cutting forces generally correlate with reduced tool wear [7].

One critical concern is the widespread use of cobalt in tool matrices. Cobalt’s toxicity [8], environmental impact [9], and price volatility have prompted extensive efforts to identify more sustainable alternatives. Alloys based on Fe, Cu, and Ti have shown promise [4,10,11,12], with mechanical properties comparable to cobalt-based systems [13]. However, certain alloyed powders remain cost-prohibitive, with prices up to 70% higher than Co [14].

Manufacturing methods also differ depending on tool geometry. Due to their complex shape, diamond beads are often produced by free (pressureless) sintering (FS), which requires slow heating rates (e.g., 2 °C min⁻¹) and long dwell times [15]. In contrast, DIT segments are typically fabricated via HP, which enhances densification, lowers sintering temperatures, and limits grain growth. Although HP is effective, it is energy-intensive and raises concerns over cost and environmental impact [16,17].

Microwave-assisted (MW) sintering has emerged as a promising alternative. By using electromagnetic radiation to induce volumetric heating, MW sintering enables faster heating, shorter cycle times, and reduced thermal gradients compared to conventional techniques [18]. While bulk metals generally reflect microwave energy due to their high electrical conductivity, fine metal powders can absorb it effectively, enabling successful sintering [19].

Despite these advantages, systematic investigation of MW sintering for DIT segments remains lacking, particularly for Fe–Cu-based matrices that are being explored as sustainable cobalt alternatives. The influence of microwave sintering on the densification behavior, microstructure, and mechanical properties of such complex composite segments is still poorly understood.

Previous studies have primarily addressed bulk alloys or beads; for instance, Wang et al. [20] achieved full densification of a Fe–Cu–Co-based alloy at 850 °C under nitrogen in 60 min using microwave sintering. Optimization with a SiC susceptor further reduced sintering time to 30 min at 900 °C [21]. However, challenges persist regarding microwave energy interactions with metallic powders and the effects of multimode cavity field distributions on uniform heating and densification [22,23].

Hybrid techniques, such as microwave hot pressing (MW–HP), which combine pressure with microwave heating, have shown promise for improved densification and shorter cycle times, as exemplified by Fe–Cu–W–Sn segments sintered at 850 °C in under 30 min [24]. Short-time sintering techniques, including Spark Plasma Sintering (SPS) and Spark Plasma Consolidation (SPC), similarly reduce diamond graphitization and limit degradation of embedded diamonds, with titanium-coated diamonds demonstrating partial protection depending on coating material and thickness [25,26]. SPS and SPC inhibit graphite formation, with SPC completely preventing catalytic carbon reactions, allowing the use of more reactive metal matrices while shortening processing time [25].

Rapid or volumetric heating techniques also improve energy efficiency. For example, low-carbon, uniform-temperature sintering of drill bits reduced energy consumption from 15.2 kWh to 6.6 kWh per segment (≈56%) while improving thermal uniformity (ΔT reduced from ~48 °C to 8–11 °C) and increasing energy utilization from 30% to 94% [27]. Analogously, microwave (MW) sintering can minimize diamond damage, allow flexible binder selection, and offer a scalable, energy-efficient alternative to SPS/SPC and conventional hot pressing. Based on experimental heating conditions used in this study, a single diamond tool segment (5.2 g) consumes ~0.167 kWh via MW (1 kW, 20 min, 2 segments) versus ~0.97 kWh per segment for hot pressing (70 kW, 10 min, 12 segments), representing a roughly six-fold reduction in energy consumption per segment.

While microwave sintering has demonstrated potential for rapid and energy-efficient densification of metal powders, its specific application to Fe–Cu-based diamond-impregnated tool (DIT) segments remains largely unexplored. The explicit research problem addressed in this study is to evaluate the feasibility and effectiveness of microwave-hybrid sintering as an energy-efficient alternative to conventional hot pressing for manufacturing DIT segments. This includes understanding how microwave processing affects densification behavior, microstructural characteristics, and mechanical performance compared to traditional methods. To date, microwave sintering has been investigated primarily for bulk matrix alloys and beads, but not systematically for DIT segments.

This study aims to address this gap by:

• Investigating the feasibility of microwave-assisted sintering for fabricating Fe–Cu-based DIT segments.
• Evaluating the densification behavior, microstructure, and mechanical performance of the microwave-sintered segments.
• Comparing these results with conventionally sintered (HP) counterparts.

By doing so, this research seeks to determine whether microwave sintering offers a viable, energy-efficient, and environmentally friendly alternative to traditional hot pressing in the manufacturing of diamond-impregnated tool segments.

2. Materials and Methods

2.1. Matrix and MMC Compacts

The composition of the matrix alloy used in this study is an industrial formulation developed by DIAPOR—Diamantes de Portugal, S.A., Rio Meão, Portugal.

Matrix powders, including cobalite CNF (d₅₀ = 11 μm), ultrafine cobalt (d₅₀ = 6 μm), iron phosphate (d₅₀ = 6 μm), tin bronze (d₅₀ = 36 μm), and tungsten carbide (d₅₀ = 5 μm), were homogenized using a Argar Mix 10 Turbula^® mixer (Willy A. Bachofen AG, Muttenz, Switzerland). For the experiments, MMC green bodies containing 5 vol.% synthetic diamond particles (d₅₀ = 360 μm) were prepared. The mixed powders were uniaxially compacted into uniformly shaped specimens (23 mm × 10 mm) using a steel mold and an Arga VCP30 hydraulic press (AFT S.p.A., Piacenza, Italy), which provides a maximum force of 300 kN. Pressing was carried out at three different pressures: 160 MPa, 350 MPa, and 610 MPa. The corresponding green densities achieved were 57%, 67%, and 77% of the HP segment density, as determined by helium gas pycnometry (Accupyc 1330, Micromeritics Instruments Corp., Norcross, GA, USA).

2.2. Sintering Experiments

Sintering cycles were conducted at Diapor S.A. using a Vulcan 70 SV hot press (70 kW), manufactured by Idea S.p.A. (Piacenza, Italy), under a controlled atmosphere consisting of Ar with 5 vol.% H₂ (99.999% purity, supplied by Sociedade Portuguesa do Ar Líquido, Arlíquido Lda., Algés, Portugal). Both the matrix and MMC segments were sintered at 850 °C with a dwell time of 3 min, using a heating rate of about 80 °C min⁻¹. The complete cycle (including cooling down to 300 °C) lasted about 18 min.

Microwave (MW) sintering was carried out using a BP–111 multimode microwave laboratory oven (2.45 GHz, up to 1 kW output), supplied by Microwave Research & Applications, Inc. (Carol Stream, IL, USA), at LNEG (Figure 1a). Temperature measurements were performed using a Pt-sheathed Type-S thermocouple positioned just above the top surface of the segment. This thermocouple was selected due to its high thermal stability and relatively low interaction with microwave radiation. It was carefully placed outside the region of maximum field intensity to minimize potential perturbation effects, such as the antenna effect reported in some studies involving metallic probes. No anomalous heating or irregularities in temperature readings were observed during the experiments. Optical pyrometry was not feasible due to the design of the microwave oven. Although fiber-optic sensors are established for temperature monitoring in microwave ovens, and low-melting-point markers could, in principle, provide surface-to-core thermal information, such measurements were beyond the scope of this study, and the marker approach remains unvalidated for microwave heating of metal-matrix composites.

The experimental setup used for MW sintering is shown in Figure 1. To minimize oxidation during sintering, the green segments (inset in Figure 1b) were placed inside a quartz tube (inner diameter: 17 mm; length: 130 mm; Figure 1c), which was positioned on a ceramic fiber support (Figure 1c). This tube was flanked by two SiC susceptors (Figure 1b) and encased in an aluminosilicate fiber pod (Figure 1d). Susceptors are materials characterized by high dielectric losses even at low temperatures, such as silicon carbide (SiC) and graphite. These materials can efficiently absorb microwave radiation at room temperature, rapidly converting it into heat. This heat is then transferred to nearby samples primarily through conduction and radiation. Initially, many materials exhibit low dielectric losses, making them poor absorbers of MW energy. However, once heated by the susceptor, their temperature increases, and their dielectric properties improve. As a result, the sample begins to couple directly with the MW field, leading to more efficient self-heating [28,29].

The sintering atmosphere was established by continuously flowing a gas mixture of argon (Ar) and 7 vol.% hydrogen (H₂), both 99.999% pure, supplied by Sociedade Portuguesa do Ar Líquido (Arlíquido Lda., Algés, Portugal), at a flow rate of 2 L min⁻¹. The gas was introduced through the thermocouple inlet (Figure 1a). To prevent overheating and protect the electronic components of the apparatus, a cooling fan was positioned on the side opposite the magnetron to promote heat dissipation. However, this configuration permitted limited ingress of ambient air into the oven, creating a potential risk of oxidation for oxygen-sensitive materials. To mitigate this effect, the samples were placed inside a quartz tube—transparent to microwaves—and exposed to the same gas mixture containing 7 vol.% H₂. This hydrogen content was sufficiently high to maintain a strongly reducing environment, yet below the flammability threshold under the experimental conditions. Consequently, oxidation was largely restricted to superficial layers, which were subsequently removed by mechanical polishing.

Although the two atmospheres employed are not strictly identical, both are strongly reducing and lie within the typical hydrogen concentration range used for the sintering of metal-matrix composites. From a thermodynamic standpoint, the effective oxygen partial pressure (pO₂) governed by the H₂–H₂O equilibrium remains exceedingly low (≈10⁻²⁰–10⁻¹⁶ atm at 800–1000 °C) for both 5% and 7% H₂. These conditions correspond to a very low dew point and minimal water activity within the gas phase, ensuring that oxidation during sintering is effectively suppressed. The marginal increase in hydrogen content only slightly reduces pO₂ and, therefore, exerts a negligible influence on densification behavior, microstructural evolution, or the mechanical performance of the sintered components, aside from minor surface oxidation.

Microwave sintering was conducted at temperatures ranging from 720 °C to 920 °C, with dwell times between 3 and 60 min, using a heating rate of ≈67 °C min⁻¹. Typically, it took around 12 min to reach 820 °C from room temperature, and the entire process—including natural cooling to 300 °C—lasted about 36 min (i.e., double the HP process), owing to the excellent insulation of the ceramic pod.

Conventional free sintering (FS) cycles were carried out in an electric horizontal tubular furnace (6.5 kW), manufactured by Termolab—Fornos Eléctricos, Lda. (Águeda, Portugal) at LNEG, for process insight and not intended for mechanical property evaluation. The sintering was performed under a controlled atmosphere of high-purity argon (99.9%), supplied by Arlíquido Lda. Temperature monitoring was achieved using a Type–K thermocouple positioned near the segments inside the alumina tube chamber. To minimize surface oxidation during sintering, the segments were placed between graphite plates. Matrix segments were conventionally sintered at temperatures between 750 °C and 950 °C, using a heating rate of 10 °C min⁻¹ and a dwell time of 5 min, under a continuous argon flow of 5 L h⁻¹.

Figure 1. Microwave setup: (a) MW oven; (b) sheathed type–S thermocouple and SiC susceptors (inset shows the top view location of the green samples); (c) quartz tube; (d) ceramic fiber pod [30].

Figure 1. Microwave setup: (a) MW oven; (b) sheathed type–S thermocouple and SiC susceptors (inset shows the top view location of the green samples); (c) quartz tube; (d) ceramic fiber pod [30].

2.3. Physical and Mechanical Characterization

The surfaces of MMC sintered segments were cleaned using a wire brush disk, while matrix segments underwent standard metallographic preparation, including grinding and polishing to a final grit size of 0.5 µm. Microstructural analysis was performed on etched samples using a FEG XL30 scanning electron microscope (Philips, Eindhoven, The Netherlands) equipped with an energy-dispersive spectrometer (EDS). Etching was carried out using a 2% Nital solution for 30 s to reveal the matrix microstructure.

Density (ρ) measurements for both MW and FS segments were performed applying the Archimedes method using a Mettler Toledo AG204 (Mettler-Toledo GmbH, Greifensee, Switzerland) analytical balance, with a 220 g maximum capacity and 0.1 mg readability. For HP segments, real density (ρᵣ) was determined via helium gas pycnometry (see Section 2.1). The porosity (ε) of the sintered segments was then estimated considering that:

ε = 1 − ρ/ρᵣ,

(1)

Dynamic Young’s modulus (E) and shear modulus (G) were determined using the RFDA System 23 (IMCE NV, Genk, Belgium), operating in flexural and torsional vibration modes, respectively, in accordance with ASTM C1259–96 [31].

Vickers hardness (HV) was measured with a 432–SVD hardness tester (Wolpert Prüftechnik BV, Aachen, Germany) applying a 1 kg for 10 s, following ISO 6507–1 [32].

Flexural strength (σᵣ) was evaluated by three-point bending tests (with a distance between external loading points of 15 mm) according to ISO 3327 [33], using an Instron 3369 universal testing machine (Norwood, MA, USA) at a constant crosshead speed of 1.0 mm min⁻¹. To ensure valid comparison under ISO 3327, all measurements were performed using a span-to-depth ratio of at least 10:1, as specified by the standard.

Attention is drawn to the fact that the MW and HP specimens differ in geometry (MW ≈ 8 mm × 4 mm × 18 mm; HP ≈ 10 mm × 3 mm × 23 mm), which may influence defect sensitivity and stress distribution, even though both comply with the ISO 3327 span-to-height criteria. Future work could harmonize specimen dimensions to enable a more direct and quantitative comparison between MW and HP sintered segments.

The jig rollers were made of WC and had a diameter of 5 mm, in accordance with the referenced standard [33]. The standard also stated that the value of $\mathbf{\Delta }\mathit{\sigma }/\mathbf{\Delta }\mathit{t}$ should not exceed 200 MPa s⁻¹. In the tests of the matrix segments produced by both HP and MW, the $\mathbf{\Delta }\mathit{\sigma }/\mathbf{\Delta }\mathit{t}$ value was on average 85 MPa s⁻¹, at a crosshead speed of 1.0 mm min⁻¹. The flexural rupture stress (${\mathit{\sigma }}_{\mathit{r}}$) was determined using the equation:

$${\sigma }_r={\displaystyle \frac{3Fl}{2b{h}^2}}$$

(2)

where $\mathit{F}$ is the fracture load (N), $\mathit{l}$ is the span between the support rollers (mm), $\mathit{b}$ is the specimen width (mm), and $\mathit{h}$ is the specimen thickness (mm).

The Weibull parameters—modulus m and characteristic strength σ₀—were estimated using the Maximum Likelihood Method in R (version 4.5.0, The R Foundation for Statistical Computing, 2025), in accordance with ASTM C1239 [34].

Fractographic examination of fractured specimens was conducted using an Olympus SZH stereomicroscope and a Philips FEG XL30 SEM. Fracture origin dimensions (depth and width) were measured using AnalySIS image analysis software (version AnalySIS FIVE). Fracture toughness (K_IC) was estimated via fractographic methods in accordance with ASTM C1322 [35]. The equation used was:

$$K_{IC}=Y{\sigma }_r\sqrt{a}$$

(3)

where ${\mathit{\sigma }}_{\mathit{r}}$ is the rupture stress, $\mathit{a}$ is the depth of the critical defect, and $\mathit{Y}$ is a material-dependent geometric factor (ranging from 1.13 to 1.99). When the value of the $\mathit{c}\mathbf{/}\mathit{a}$ ratio (where $\mathit{c}$ is the width of the critical defect) fell between tabulated intervals, the corresponding $\mathit{Y}$ value was obtained by linear interpolation. The uncertainty in the estimated ${\mathit{K}}_{\mathit{I}\mathit{C}}$ values arises primarily from the accuracy of defect size measurements and the identification of the fracture origin, and is typically within ±10–15%.

3. Results and Discussion

3.1. Optimization of MW Thermal Cycle

Figure 2 presents the densification curves for matrix segments sintered via MW and conventional FS, using compacts with an initial green density of 57%, 67% and 77%.

All curves exhibit similar overall profiles; however, the MW curve is noticeably shifted toward lower temperatures—by roughly 100 °C—compared to the FS curve, as shown in Figure 2a. Additionally, the MW sintering curve shows a significantly broader plateau. For example, at 765 °C, the MW-sintered segments achieved a relative density of 93%, whereas those processed by FS reached only 63%.

Under the assumption that the optimum sintering temperature corresponds to achieving 95% of the theoretical density, the sintering plateau for the MW samples spans a range from 800 °C to 870 °C. In contrast, the FS samples reach this densification level at around 920 °C. These results indicate that, for segments with an initial green density of 57%, the application of microwave sintering leads to a reduction in the effective sintering temperature by circa 90 °C, highlighting the enhanced densification efficiency associated with microwave heating.

Although not shown in Figure 2a (for clarity), the densification curves for matrix segments with a green density of 67% revealed that the optimum sintering temperature for the MW–67% series was around 770 °C, corresponding to a relative density of 92%. At the same temperature, the FS–67% series achieved only 75% density. Maximum densification for the FS–67% samples occurred at 920 °C, reaching a relative density of 95%. Similarly, for the MW–77% series, optimum densification (92–93%) was obtained within the temperature range of 720–750 °C. In contrast, the FS–77% samples only reached similar densities at 920 °C, with a maximum of 94%.

Figure 2b illustrates the effect of green density on the sintering behavior of segments processed by microwave (MW). Up to 750 °C, a direct relationship is observed: the higher the green density, the higher the final sintered density. However, above 750 °C, two distinct behaviors emerge. In the MW–57% series, densification continued to increase with temperature. Conversely, in the MW–67% and MW–77% series, a marginal decrease in final density was observed beyond 800 °C and 770 °C, respectively. Notably, this reduction in density is not attributed to increased porosity, but rather to non-uniform volumetric expansion during sintering. A similar phenomenon was described by Takayama et al. [36], who noted that, as the temperature rises, particle neck growth accelerates due to enhanced diffusion mechanisms. Beyond a critical threshold, the compact begins to behave as a bulk metal, causing a significant reduction in microwave penetration. This occurs because the skin effect limits electromagnetic wave absorption to the material’s surface, effectively shielding the interior from further microwave exposure. As a result, densification occurs mainly at the surface, while the core remains under-sintered. In some cases, this non-uniform densification can induce surface distortion or cracking due to internal stress gradients. This effect was not observed in the MW–57% series. A plausible explanation lies in the higher porosity of the green body, which results in lower thermal conductivity and thus allows deeper microwave penetration. This promotes more uniform volumetric heating and densification [37,38].

Regarding the optimum sintering temperature:

▪

MW–77%: between 720 °C (92%) and 750 °C (93%);

▪

MW–67%: between 720 °C and 770 °C (both at 92%);

▪

MW–57%: widest sintering plateau, ranging from 800 °C to 870 °C with relative densities above 94%.

These results suggest that increasing green density lowers the optimum sintering temperature. However, it also narrows the sintering plateau, potentially making process control more challenging due to the shielding effect observed in higher-density compacts.

For the MW–57% series, relative densities ranging from 93% to 95% were observed within the 770 °C to 850 °C sintering range, indicating the presence of closed porosity. Such porosity can be eliminated either by applying external pressure or by partial melting of a matrix component, which then fills the pores via capillary action. At 870 °C, the density increased to 97%, indicating the onset of melting in one such component, most likely tin bronze. Tin bronze (Cu–12% Sn) has a melting point near 900 °C, supporting the hypothesis that partial melting begins around this temperature. However, sintering at 920 °C resulted in segment deformation, likely due to excessive liquid-phase formation, which compromises structural integrity. This also explains the decrease in density observed in Figure 2a for the MW–57% series at 920 °C, along with the increased data scatter. Given that the most promising results were obtained for the MW–57% samples, the effect of dwell time on their densification behavior was investigated at three sintering temperatures: 720 °C, 820 °C, and 870 °C. The results are shown in Figure 3.

At 820 °C, the densification curve shows minimal variation in density with increasing dwell time. Segments reached approximately 95% relative density after just 3 min of sintering, with no significant improvement observed at longer dwell times. This suggests that at this temperature, densification is effectively completed early in the process. Additionally, the absence of significant liquid-phase formation at 820 °C may explain the stabilization of density over time.

At 870 °C, a 2% increase in average density is observed from 97% at 5 min to 99% at 30 min. Since the segments already exceed 95% density at this stage, the further increase may indicate the onset of liquid-phase sintering, likely due to partial melting of a matrix constituent (e.g., tin bronze). This densification behavior is consistent with capillary-driven liquid-phase mechanisms. In contrast, the curve for 720 °C exhibits non-linear behavior, with noticeably higher standard deviations (above 3%) for dwell times of 5, 15, and 30 min. This variation suggests unstable or incomplete densification, likely because 720 °C represents an intermediate sintering temperature where the diffusion-driven mechanisms are insufficiently activated. Thus, dwell time does not appear to significantly influence densification at this temperature.

As previously discussed, increasing the green density results in a reduction in the required sintering temperature; however, this also narrows the sintering plateau. Notably, the MW–57% series yielded the highest overall density values (≈95%). Therefore, it is concluded that for the selected matrix composition, green density should not exceed 57%.

Densification analysis for the MW–57% series identified a sintering plateau extending from 800 °C to 870 °C. Based on this, an intermediate sintering temperature of 820 °C was selected to provide a buffer against minor temperature fluctuations during processing. Although no significant dependency between dwell time and final density was found, a dwell time of 5 min was chosen for microwave sintering to improve process stability and temperature control.

3.2. Physical and Mechanical Properties of HP and MW Sintered Segments

Densities (ρ), porosity (ε), dynamic Young’s modulus (E), shear modulus (G), flexural strength (${\mathbf{\sigma }}_{\mathbf{r}}$), fracture toughness (K_IC), and Vickers hardness (HV1) values for the matrix and MMC segments produced by HP and MW are summarized in

Table 1. Densities, porosity (ε), dynamic Young’s modulus (E), shear modulus (G), flexural strength (${\mathbf{\sigma }}_{\mathbf{r}}$), fracture toughness (K_IC) and Vickers hardness (HV1) values for the matrix and MMC segments produced by MW sintering (MW matrix; MW–MMC) and hot pressing (HP matrix; HP–MMC).

Series	n	ρ (Mg m−3)	Rel.Dens. (%)	ε (%)	E (GPa)	G (GPa)	σr (MPa)	KIC $(\mathbf{M}\mathbf{P}\mathbf{a} {\mathbf{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.})$	HV1
MW–Matrix	15	$7.78\pm 0.07$	$94.4\pm 0.1$	$5-6$	$165\pm 3$	$64\pm 3$	1106$\pm$120	14.0$\pm$1.5	$260\pm 37$
HP–Matrix	12	$8.24\pm 0.01$	$100$	$-$	$193\pm 2$	$76\pm 1$	1246$\pm$78	14.4$\pm$1.9	$331\pm 24$
MW–MMC	14	$7.80\pm 0.07$	$96.2\pm 0.1$	$4-5$	$178\pm 5$	$68\pm 3$	1039$\pm$83	14.8$\pm$1.7	$-$
HP–MMC	13	$8.11\pm 0.01$	$100$	$-$	$196\pm 3$	$75\pm 1$	986$\pm$120	15.8$\pm$1.5	$-$

.

MW segments were sintered under the optimized conditions outlined in Section 3.1, resulting in final dimensions of circa 8 mm × 4 mm × 18 mm (width × height × length), compared to 10 mm × 3 mm × 23 mm for the HP counterparts. Based on the density values obtained, the estimated porosity in MW segments was approximately 5–6% for the matrix segments and 4–5% for MMC segments. As previously discussed, these values are associated with closed porosity, which is difficult to eliminate in the absence of applied pressure. No significant variations in density were found between different MW production cycles for either the matrix or MMC segments. Specifically, densities of 7.78 ± 0.07 Mg m⁻³ and 7.80 ± 0.06 Mg m⁻³ were measured for the MO matrix and MO–MMC segments, respectively. The standard deviation of less than 1% confirms the reproducibility of the MW sintering process. The relative density values for the MW–MMC series (96.2 ± 0.1%) fall within the range reported by Schmidt et al. [39] for Co-based powders containing diamond particles sintered by MW (95–97%).

The HP series exhibited higher E and G values—by circa 10–15%—compared to the MW series. No significant differences were observed among the samples tested within the HP series. However, in the MW series, the MW–MMC samples showed higher E and G values (about 6–8%) compared to the MW matrix samples. Since E and G are intrinsic material properties dependent on chemical composition, microstructure, and the presence of defects (e.g., pores or particle agglomerates), the higher porosity in MW segments likely explains their comparatively lower values [40].

Regarding Vickers’ hardness, the MW matrix segments showed lower mean values (260 ± 37 HV1) than those of the HP matrix segments (331 ± 24 HV1). As hardness reflects a material’s resistance to deformation, indentation, or penetration, this reduction can be attributed to the increased porosity in MW segments. These findings are consistent with those reported by Cygan-Bączek et al. [14], who observed values of 299 ± 7 HV1 for a Fe–Mn–Cu–Sn–C base material sintered by spark plasma sintering (SPS) at 900 °C for 10 min under 35 MPa. Similar HV1 values (200–260) were also reported for Fe–P–Ni–Cu–Sn alloys processed by FS at 840–925 °C [41].

A clear trend was observed, with higher flexural strength (${\sigma }_r$) values measured for the MW–Matrix (1106 ± 120 MPa) and HP–Matrix (1246 ± 78 MPa) compared to the MW–MMC (1039 ± 83 MPa) and HP–MMC (986 ± 120 MPa) series. This difference is likely attributed to the size of the fracture origin. In the MMC segments, the fracture origin is typically associated with diamond particles located on the surface or sub-surface, with dimensions of around 400 μm. In contrast, in the matrix segments, the fracture origin is smaller, typically ranging between 100 μm and 200 μm.

Based on the results of the three-point bending tests, the Weibull modulus ($\widehat{m}$_U) and characteristic strength (σ₀) were determined using the Weibull statistical distribution. The values obtained are summarized in

Table 2. Estimated Weibull parameters (with 90% confidence bounds).

Serie	$\widehat{\mathit{m}}$U	σ0 (MPa)	$\widehat{\mathit{m}}$upper	$\widehat{\mathit{m}}$lower	σ0,upper (MPa)	σ0,lower (MPa)
MW matrix	11.3	1157	13.8	8.4	1193	1123
HP matrix	16.8	1283	20.5	12.6	1309	1258
MW–MMC	16.4	1074	20.5	11.5	1101	1048
HP–MMC	9.4	1058	11.8	6.6	1105	1014

. Weibull modulus values ranged from 9.4 (HP–MMC) to 16.8 (HP matrix), without a clear dependence on the sintering method. The observed variation is likely associated with the size and distribution of fracture origins and may be minimized by increasing the number of tested segments (n = 12–15 in this study).

The Weibull modulus characterizes the distribution of a material’s strength. The higher the value of $\widehat{m}$_U, the more homogeneous the material, meaning that critical defects are uniformly distributed throughout the volume and that the probability density function curve of strength will be narrower [34]. The m value was higher for the HP matrix and MW–MMC segments, indicating a more homogeneous distribution of defects in these samples. In contrast, the lower $\widehat{m}$_U values observed in the other segments are likely attributable to the agglomeration of diamond particles within the matrix. These agglomerates can act as stress concentrators, promoting premature failure at lower rupture strengths and contributing to variability in the measured data, as illustrated in Figure 4. Although a Turbula^® mixer was employed to ensure uniform blending, its low shear intensity may not be sufficient to fully break up diamond clusters. Further optimization of the mixing process—such as adjusting mixing time, incorporating a binder, or pre-deagglomerating the diamond particles—could enhance dispersion and improve microstructural uniformity.

It should be noted that the mixing process followed a fixed industrial protocol, and modifications to the parameters were not permitted for comparison purposes. While the microstructural variability is qualitatively linked to residual agglomeration, no quantitative dispersion or stereological analysis was performed. This is acknowledged as a limitation and a direction for future work.

For the MMC series, the measured ${\mathsf{\sigma }}_{\mathrm{r}}$ values were similar, as the fracture origin was consistently associated with diamond particles. In contrast, for the matrix segments, SEM–EDS analysis revealed that the fracture origin was linked to tin bronze agglomerates (as shown in Figure 5a).

The variation in ${\mathsf{\sigma }}_{\mathrm{r}}$values may therefore be attributed to differences in microstructural homogeneity. The HP matrix series appears more homogeneous, potentially resulting in a more uniform distribution of tin bronze particles (Figure 5b).

Scarce data exist on the fracture behavior of MW-sintered parts. Peng et al. [42] investigated the mechanical properties and microstructure of MW- and FS-sintered Fe–2Cu–0.6C alloys at temperatures up to 1150 °C. They concluded that fracture in MW-sintered parts was closely related to the resulting microstructural features, exhibiting a mixed fracture mode involving both ductile and brittle mechanisms, in contrast to the predominantly brittle fracture observed in conventionally sintered counterparts.

In the present work, brittle fracture was the predominant failure mechanism in the MW-sintered matrix, as illustrated in Figure 6.

To improve the fracture toughness of diamond tool segments, it is of utmost importance to accurately investigate the underlying failure mechanisms and establish the structure–property relationships that link fracture initiation to the microstructure. Griffith was the first to define such a relationship by establishing a fundamental connection between material strength, toughness, and the size of critical flaws [43]. According to the data presented in Table 1, the differences in the estimated values of fracture toughness (K_IC) were not statistically significant—circa 14 ± $2 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$for the matrix and 15 ± $2 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$ for the MMC segments—regardless of the manufacturing process used. These values are typical of brittle materials such as cemented carbides, ceramics, tool composites, and metal matrix composites, all of which generally exhibit markedly lower fracture toughness than metals [44]. For example, the fracture toughness (K_IC) of ceramics—such as alumina, zirconia, and silicon nitride—typically ranges from 2 to $7 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$, depending on composition, grain size, and porosity [45]. Although certain toughened ceramics (e.g., transformation-toughened zirconia) may exceed $10 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$, their toughness remains lower than that of most metals [46].

In contrast, cemented carbides—widely used in cutting tools—can exhibit higher K_IC values, typically ranging from 8 to $30 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$, depending on the tungsten carbide (WC) particle size and the content of the metallic binder phase [44]. An increase in the mean free path of the binder phase promotes plastic deformation and crack bridging, which enhances toughness [47].

The fracture toughness values measured in this study (~14–$15 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$) are thus well aligned with the performance expected of hard, brittle composite materials. They are higher than most ceramics and comparable to medium-toughness cemented carbides and metal matrix composites engineered for tool applications.

The results of this study demonstrate that not only does microwave sintering achieve mechanical properties comparable to conventional hot pressing, it does so at significantly reduced sintering temperatures and shorter processing times, without the need for external pressure. This represents a notable advancement in the manufacturing of diamond tool composites, offering a more energy-efficient and potentially scalable alternative to traditional sintering methods. Additionally, the comprehensive analysis of densification behavior and microstructural evolution under microwave heating enhances the fundamental understanding of microwave–material interactions in complex metal matrix composites. These insights expand the practical applicability of microwave sintering in industrial contexts, moving beyond proof-of-concept demonstrations to a validated processing route for high-performance tool materials. Additionally, this study systematically mapped the densification behavior of Fe–Cu-based DIT segments as a function of temperature, green density, and dwell time, identifying 820 °C, 5 min, and 57% green density as the optimal microwave sintering condition; the observed nonlinear densification trends were explained through established mechanisms such as the skin effect and capillary-driven liquid phase sintering. Attention is drawn to the fact that the discussion of skin effects at higher green densities is based solely on bulk Archimedes density measurements. The observed variations in overall density suggest that surface-to-core densification differences may exist; however, these measurements provide no spatially resolved information to confirm this. Techniques such as microhardness mapping, quantitative metallography, or µCT imaging are required to determine whether local density or porosity gradients are present. Future work should incorporate these methods to provide a more detailed understanding of local densification behavior in MW-sintered diamond tool segments.

Previous studies on microwave (MW) sintering of metallic materials have largely been limited to academic or exploratory investigations, often involving idealized systems such as pure Fe–Cu or Fe–C alloys. In contrast, the present work applies MW sintering to a complex, industrial-grade matrix (Co–Fe–Sn–bronze–WC) incorporating abrasive diamond particles, thereby replicating real tool segment production conditions. This demonstrates not only the scalability of the process but also its industrial relevance. Preliminary cutting tests on granite confirmed the excellent performance of the MW-sintered segments, further supporting the practical applicability of the approach; detailed results will be presented in future work.

4. Conclusions

Microwave (MW) sintering significantly enhanced the densification of diamond tool segments, enabling a reduction in optimum sintering temperature by 90–170 °C compared to conventional free sintering (FS), depending on the initial green density. Increasing the green density from 57% to 77% lowered the required temperature by ~80 °C and narrowed the sintering plateau from 75 °C to 25 °C. The highest relative density (~95%) was achieved at 57% green density.

Under optimal MW conditions—57% green density, 820 °C, and 5 min dwell time—near-full densification was obtained. Matrix segments exhibited ~15% lower Young’s modulus and ~20% lower Vickers hardness than hot-pressed (HP) counterparts, while MMC segments showed ~10% lower modulus, likely due to residual porosity in the absence of applied pressure. Flexural strength of MW-processed matrix segments (1106 ± 120 MPa) was slightly lower than HP (1246 ± 78 MPa), whereas MMC segments with diamond particles exhibited statistically comparable flexural strengths between MW (1039 ± 83 MPa) and HP (986 ± 120 MPa). The Weibull modulus ranged from 9 to 17, reflecting flaw variability, and fracture toughness remained consistent across all samples (14 ± 2 to 15 ± $2 \mathrm{MPa} {\mathrm{m}}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$).

Overall, MW sintering under optimized conditions produced MMC segments with mechanical performance comparable to HP, demonstrating its potential as a pressureless, energy-efficient alternative for high-performance diamond tool manufacturing. Energy consumption per segment was reduced roughly six-fold (0.167 kWh via MW versus 0.97 kWh per segment for HP).

This study extends previous work by applying MW sintering to complex, industrial-grade Fe–Cu-based composites, establishing optimized processing parameters, and revealing key process–structure–property relationships. Insights into microwave–material interactions, particularly the influence of porosity on energy absorption and densification, further support the scalability of this approach.

By validating MW sintering as a fast, sustainable, and scalable route for fabricating diamond-impregnated composites, this work advances the application of microwave technology in industrial abrasive tool production.