Graphical abstract

The addition of 5 wt.% nano-Al₂O₃ alters the sintering behavior of copper under SPS. The nanoparticles inhibit grain growth through boundary pinning, resulting in refined grains, enhanced electrical conductivity, but reduced hardness compared with fully densified coarse-grained pure copper.

1. Introduction

Copper (Cu) is widely used in electrical and thermal applications due to its excellent electrical and thermal conductivity and high ductility [1-5]. However, the relatively low mechanical strength and poor high-temperature stability of pure Cu limit its use in demanding engineering applications where high strength and wear resistance are required [6-8]. Reinforcement of Cu with ceramic dispersoids (Al₂O₃, ZrO₂, SiO₂, etc.) is an established route to enhance high-temperature strength and wear resistance while retaining reasonable conductivity [9-14]. Dispersion-strengthened Cu exhibits improved high-temperature performance due to particle pinning of grain and sub-grain boundaries and Orowan strengthening mechanisms.

Alumina (Al₂O₃) is an attractive reinforcement for Cu due to its high melting point, high elastic modulus and hardness, chemical stability, and good wear resistance [15-18]. When incorporated into a Cu matrix, Al₂O₃ particles can significantly increase strength and wear resistance; these properties make Cu-Al₂O₃ composites relevant for industrial applications such as spot-welding electrodes, heat-dissipating components, and wear-resistant electrical contacts.

Conventional consolidated processes for Cu-Al₂O₃ composites include pressure-less sintering, hot pressing (HP), hot isostatic pressing (HIP), hot extrusion, and infiltration [19-23]. These methods suffer from various limitations, such as high sintering temperatures and long cycle times, leading to grain growth and coarsening of dispersoids, poor interfacial bonding or agglomeration of ceramic particles, and difficulty in achieving full densification without excessive matrix grain growth. Such shortcomings can degrade mechanical and functional properties of composites.

Spark plasma sintering (SPS), also referred to as pulsed electric current sintering (PECS), has emerged as an effective consolidation technique for metal matrix composites because it enables very high heating rates, short dwell times, and relatively low sintering temperatures, promoting rapid densification while suppressing matrix grain growth and limiting particle coarsening [24-29]. These characteristics often result in finer microstructures, improved particle-matrix interfaces, and enhanced mechanical and functional properties compared with conventional sintering approaches. SPS has therefore been widely applied to Cu-Al₂O₃ and related systems. However, most previous studies primarily focus on composite samples alone, while the intrinsic sintering response of pure Cu processed under identical SPS conditions is rarely examined as a direct benchmark. As a result, the specific role of Al₂O₃ addition in modifying densification behavior, grain evolution, and transport properties during SPS remains insufficiently clarified.

The present work aims to compare the microstructure and properties of pure Cu and a Cu-5 wt.% Al₂O₃ composite consolidated by SPS across a range of sintering temperatures. The objectives are to evaluate densification and grain/particle evolution with temperature, characterize particle-matrix interfaces, and correlate microstructural features with hardness and electrical performance to identify optimal sintering conditions for balanced properties in Cu-Al₂O₃ composites.

2. Experimental Section

2.1 Materials and Methods

Gas-atomized copper powder (99% purity, particle size 15-45 μm) and nano-sized Al₂O₃ powder (99% purity, average particle size of 135 nm) were used as the starting materials. The powders were weighed to obtain a composition of 95 wt.% Cu and 5 wt.% Al₂O₃, followed by mixing using a laboratory-scale horizontal roller ball mill at 200 rpm for 4 h. A reinforcement content of 5 wt.% Al₂O₃ was intentionally selected to promote strong grain-boundary pinning and the formation of Al₂O₃-rich boundary regions, enabling a clear assessment of microstructural evolution and property trade-offs under identical SPS conditions. The morphology of the starting powders and Cu-5 wt.% Al₂O₃ powder after mixing is shown in Fig. 1. The EDS elemental mappings confirm the relatively homogeneous distribution of Al₂O₃ nanoparticles over the Cu surfaces.

Consolidation of pure Cu and Cu-5 wt.% Al₂O₃ powders was carried out using a spark plasma sintering system (SPS-515S, Dr. Sinter Lab, Japan). These powders were loaded into a graphite die with an inner diameter of 15 mm, and a uniaxial pressure of 50 MPa was applied throughout sintering process. The samples were heated at a rate of 100°C/min and sintered at three different temperatures at 800°C, 900°C, and 1000°C, and held for 5 min before rapid cooling under a water-assisted cooling system. To prevent adhesion between the mold and the compact, the graphite die and punches were coated with a thin graphite layer before sintering.

2.2 Characterization

The consolidated samples were subsequently characterized as density, microstructure, mechanical, and electrical properties. Bulk density was measured according to Archimedes’ principle. Microstructural observations were performed using optical microscopy (OM, Olympus 53M, Japan), scanning electron microscopy (SEM, JEOL JSM-6500F, Japan), electron backscatter diffraction (EBSD, FE-SEM-S500 equipped with Velocity™ EBSD camera system, Japan), and X-ray diffraction (XRD, Rigaku Ultima IV, Japan).

Mechanical performance was assessed through Vickers microhardness testing (Mitutoyo HM-220D, Japan) using a load of 25 gf and a dwell time of 10 s. Electrical conductivity measurements were conducted on polished samples at room temperature using the impedance grain phase analyzer (Hewlett Packard 4194A). To ensure reliability, all density, microhardness, and electrical conductivity measurements were repeated five times under identical conditions, and the average values were reported.

3. Results and Discussion

3.1 Density and microstructure of sintered samples

Fig. 2 illustrates the evolution of relative density for pure Cu and the Cu-5 wt.% Al₂O₃ samples as a function of sintering temperature. It can be seen that the pure copper sample exhibits consistently high densification, increasing slightly from approximately 98.6% at 800 °C to around 99.7% at 1000 °C, corresponding to a monotonic decrease in apparent porosity with increasing SPS temperature. In contrast, the Cu-5 wt.% Al₂O₃ composite samples demonstrate lower relative density across all temperatures, indicating a higher retained porosity compared with pure Cu. The composite sample shows a slight decrease in density at 900 °C, reaching a minimum of 94.6%, before increasing again to 95.6% at 1000 °C, which implies a non-monotonic evolution of porosity with sintering temperature. The reduced densification at 900 °C is attributed to the combined effects of particle rearrangement limitation and pore stabilization induced by Al₂O₃-rich regions at particle boundaries. At this intermediate temperature, the diffusion activity of Cu is enhanced compared to 800 °C but remains insufficient to overcome the diffusion barrier imposed by the oxide phase, leading to pore coalescence and partial pore entrapment rather than effective pore elimination, thereby temporarily increasing the effective porosity of the composite. When the sintering temperature is increased to 1000 °C, further enhancement of atomic diffusion and localized plastic deformation of the Cu matrix under the applied pressure promotes pore shrinkage and interparticle neck growth, resulting in a reduction of residual porosity and improved interparticle connectivity. This results in a partial breakdown of the pore-stabilizing effect of Al₂O₃-rich boundary regions and enables additional densification, although the final density remains lower than that of pure Cu due to the persistent constraint imposed by the oxide phase, which continues to limit complete pore elimination.

Fig. 3 reveals distinct differences in the microstructural evolution of pure Cu and the Cu-5 wt.% Al₂O₃ composite samples during SPS. For pure Cu samples, pore elimination is evident with increasing temperature. At 800 °C, residual pores (indicated by red arrows) are primarily localized at powder boundaries, whereas at 1000 °C, significant blurring of the powder boundaries with reduced porosity is observed. In contrast, the Cu-5 wt.% Al₂O₃ composite samples exhibit restricted grain growth and persistent porosity across all temperatures. The Al₂O₃ particles (indicated by yellow arrows) are predominantly situated at triple junctions of the copper powder boundary, exerting a pinning effect [30-32] that impedes grain boundary migration and matrix diffusion. At 1000°C, although enhanced diffusion promotes partial pore elimination, the Al₂O₃ particles continue to act as diffusion barriers, resulting in significantly higher residual porosity compared to pure copper.

For a more detailed assessment of the microstructural evolution, SEM observations of the sintered samples are presented in Fig. 4. For the pure Cu samples, pore size progressively decreases, and grain growth becomes more pronounced with increasing sintering temperature. At 900 °C, the well-defined annealing twins, typical of face-centered cubic copper, are clearly visible within the grains, while substantial grain coarsening occurs at 1000 °C. In contrast, the Cu-5 wt.% Al₂O₃ composite samples exhibit a markedly different behavior. The ceramic particles preferentially segregate along Cu grain boundaries, forming Al₂O₃-rich regions at 800 °C. At 900 °C, significant residual porosity remains concentrated near the Cu/Al₂O₃ interfaces, which explains the decrease in density at this temperature. Although higher temperatures enhance bonding between Cu grains, the ceramic phase continues to pin grain boundaries and restrict diffusion, resulting in persistent porosity and limited densification even at 1000 °C. These microstructural characteristics strongly influence the resulting material properties, as grain boundary integrity and second-phase distribution play critical roles in determining the mechanical performance of the composites.

The texture characteristics of the pure Cu and Cu-5 wt.% Al₂O₃ composite samples sintered at 1000 °C were evaluated using the EBSD maps presented in Fig. 5. Because the Al₂O₃ particles are nanometric and lie below the EBSD spatial resolution, these obtained maps primarily represent the crystallographic behavior of the Cu matrix. The results indicate that pure Cu undergoes extensive recrystallization and grain coarsening during sintering, whereas the Cu-5 wt.% Al₂O₃ composite maintains a refined grain structure with clear orientation interruptions caused by grain boundary pinning effect of the dispersed Al₂O₃ particles, which suppressed grain growth. The image quality (IQ) maps reveal that pure Cu develops a bimodal grain structure consisting of coarse recrystallized grains (35~70 μm) surrounded by much finer grains, reflecting active boundary migration at high temperature. In contrast, the composite exhibits a more homogeneous equiaxed grain structure with significantly smaller grains due to the pinning effect of the dispersed Al₂O₃ particles. Quantitative grain boundary misorientation analysis further shows that pure Cu contains a high-angle grain boundary (HAGB, >15°) fraction of approximately 26.4%, whereas the composite exhibits a markedly higher HAGB fraction of ~50.9%, accompanied by a reduced low-angle grain boundary (LAGB, 2-15°) fraction of ~9.3%. The grain boundary (GB) maps show that both samples contain predominantly HAGBs. However, the composite displays a higher HAGB fraction and markedly fewer annealing twins than pure Cu, consistent with restricted recrystallization and reduced boundary mobility. The increased HAGB fraction in the composite is attributed to oxide-pinning-induced grain subdivision and limited grain boundary migration rather than enhanced grain growth. The inverse pole figure (IPF) maps further show that pure Cu possesses large grains with a completely random orientation distribution, characteristic of unrestricted recrystallization, whereas the composite displays smaller grains whose orientations are locally disrupted by grain-boundary Al₂O₃-rich regions. In addition, the grain size distribution of these samples is illustrated in Fig. 6. The pure Cu sample exhibits an average grain size of 57.83 μm with a broad distribution ranging from approximately 7 to 170 μm, indicating extensive grain growth during sintering. In contrast, the composite sample shows a smaller average grain size of 36.54 μm and a noticeably narrower distribution of 5 to 96 μm.

Higher resolution EBSD analyses were performed to observe the structure at the grain-boundary Al₂O₃-rich region in the composite samples sintered at 1000 °C, as shown in Fig. 7. The IQ map reveals a highly heterogeneous microstructure in which Cu grains are fragmented and irregularly shaped due to the presence of densely packed ceramic particles. The phase map confirms that the Al₂O₃ phase surrounds and separates the Cu grains, effectively impeding their coalescence and restricting grain boundary migration. Quantitative phase analysis indicates that although the total Al₂O₃ fraction is ~10.3%, its partition fraction within grain-boundary regions increases to ~12.9%, demonstrating preferential Al₂O₃ enrichment at grain boundaries .This morphology indicates that although local diffusion attempts to promote bonding between adjacent Cu grains, the nano-oxide phase acts as a strong physical barrier, preventing full densification in this region. The IPF map further shows that the crystallographic orientations of the Cu grains in the Al₂O₃-rich zone are completely random, with no evidence of texture development. This lack of preferred orientation suggests that the grain-boundary Al₂O₃-rich regions interrupt recrystallization pathways and inhibit grain growth. These results demonstrate that the nano-Al₂O₃ phase plays a dominant role in controlling the local microstructure by blocking Cu diffusion and limiting grain coalescence.

Fig. 8 presents the XRD patterns of the pure Cu and Cu-5 wt.% Al₂O₃ composite samples sintered at 800, 900, and 1000 °C. For the pure Cu samples, only the characteristic diffraction peaks of face-centered cubic Cu are detected, indicating the absence of secondary phases formed during the SPS process. In the Cu-5 wt.% Al₂O₃ composites, the diffraction peaks corresponding to Al₂O₃ are significantly weaker and broader than those of the Cu matrix. Although the nominal Al₂O₃ content corresponds to a non-negligible volume fraction, the weak Al₂O₃ reflections are associated with the combined effects of the nanoscale size of the Al₂O₃ particles, their high dispersion within the Cu matrix, and the dominance of the strong Cu diffraction peaks. The nanosized Al₂O₃ particles result in pronounced peak broadening and reduced peak intensity, while partial peak overlapping with the Cu reflections further limits their detectability in conventional XRD scans. Furthermore, the diffraction patterns of all samples exhibit similar features across the different sintering temperatures. These results indicate that no significant phase transformation, impurity introduction, or interfacial reaction occurs during SPS, confirming the thermodynamic stability of both Cu and Al₂O₃ under the applied sintering conditions.

3.2 Mechanical properties and electrical conductivity of sintered samples

The hardness evolution of the pure Cu and Cu-5 wt.% Al₂O₃ composite samples after sintering is presented in Fig. 9. For pure Cu samples, hardness increases from 63 HV at 800 °C to 96 HV at 900 °C, and then slightly decreases to around 88 HV at 1000 °C due to grain coarsening effect dominant. The maximum hardness observed at 900 °C reflects an optimal balance between enhanced densification, which effectively reduces porosity (Fig. 2), and limited grain growth, allowing grain boundary strengthening to be largely retained. The composite samples display lower hardness values than pure Cu samples at all temperatures, increasing from 56 HV at 800 °C to 76 HV at 900 °C before slightly dropping at 1000 °C. The reduced hardness in the composite samples compared with pure Cu samples is attributed to the presence of grain-boundary Al₂O₃-rich regions that inhibit uniform grain growth and lead to retained porosity, limiting the strengthening contribution from grain refinement. Although grain refinement would normally enhance hardness according to the Hall-Petch relationship, the strengthening contribution from finer grains in the composite is substantially offset by the detrimental effect of residual porosity. In addition, the effective reinforcement of Al₂O₃ particles strongly depends on interfacial load transfer. The absence of interfacial reaction phases suggests that the Al₂O₃/Cu interface is primarily mechanically bonded, which limits stress transfer from the Cu matrix to the ceramic phase. Consequently, the combined effects of residual porosity and limited interfacial load transfer outweigh the potential strengthening from the hard ceramic phase and grain refinement, resulting in lower hardness compared to near-fully dense pure Cu. While grain structure evolution governs the overall temperature-dependent trend in hardness for both materials, the absolute hardness of the composite is primarily controlled by its porosity-dominated response.

Fig. 10 presents the apparent electrical conductivity of pure Cu and Cu-5 wt.% Al₂O₃ composites sintered at different temperatures. As expected, the composite samples exhibit consistently lower conductivity than pure Cu at all sintering temperatures, primarily due to the presence of the electrically insulating Al₂O₃ phase and microstructural discontinuities that enhance electron scattering. For pure Cu samples, the apparent electrical conductivity increases monotonically from approximately 1.89 MS/m at 800 °C to 2.25 MS/m at 1000 °C. This behavior indicates that, for nearly fully dense pure Cu, electrical transport is not governed solely by densification but is also influenced by microstructural recovery and improved interparticle bonding at higher sintering temperatures, which reduce defect- and interface-related electron scattering. In contrast, the Cu-5 wt.% Al₂O₃ composite samples exhibit a non-monotonic trend, with the apparent electrical conductivity decreasing from about 1.53 MS/m at 800 °C to a minimum of approximately 1.00 MS/m at 900 °C, followed by a pronounced increase to around 1.90 MS/m at 1000 °C. The reduced conductivity at 900 °C can be attributed to the combined effects of retained porosity, finer grain size, and Al₂O₃-rich regions, which collectively increase grain-boundary- and interface-related electron scattering. At 1000 °C, improved interparticle contact and partial grain coarsening mitigate these scattering effects, leading to a recovery of the apparent electrical conductivity. It should be noted that the measured values represent the apparent electrical conductivity of sintered compacts, which is influenced by porosity, interparticle contact quality, and contact resistance. Therefore, the data are mainly intended for comparative trend analysis rather than intrinsic conductivity evaluation.

4. Conclusion

In this study, the microstructure and properties of pure Cu and Cu-5 wt.% Al₂O₃ composites fabricated by SPS process at 800-1000 °C were investigated. The results indicated that the introduction of 5 wt.% nano-Al₂O₃ markedly alters the sintering response and final properties of copper. Pure Cu samples achieved near-full densification (98.6-99.7%) and developed a coarsen microstructure with an average grain size of 57.83 μm, whereas the composite exhibited lower density (94.6-95.6%) and significantly refined grains averaging 36.54 μm due to the strong boundary-pinning effect of the nano-oxide phase. Notably, the composite showed a density minimum at an intermediate sintering temperature, highlighting the competing effects between enhanced Cu diffusion and oxide-induced pore stabilization. High-resolution EBSD confirmed that the grain-boundary Al₂O₃-rich regions restrict Cu grain coalescence and produce disrupted orientation fields without inducing any reaction phases during sintering. These microstructural distinctions are reflected in the mechanical and electrical properties. The pure Cu samples reached a maximum hardness of 96 HV at 900 °C, while the composite samples remained softer (56-76 HV) primarily due to retained porosity that limits effective grain-boundary strengthening. In addition, the composite samples exhibit a non-monotonic change in apparent electrical conductivity with sintering temperature, while pure Cu consistently shows higher apparent electrical conductivity across all temperatures.

Article information

Funding

(1) These results were supported by "Regional Innovation Strategy (RIS)" through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (MOE)(2021RIS-003).

(2) This work was supported by the Technology Innovation Program (No. 20017647, Development of oxide dispersion strengthened superalloy materials and manufacturing technology for hypersonic engines.) funded by the Ministry of Trade, Industry & Energy (MOTIE, Korea).

Conflict of Interest

The authors have no conflicts of interest to declare.

Data Availability Statement

All data generated or analyzed during this study are includedin this article.

Author Information and Contribution

Dinh Van Cong: PhD candidate; responsible for conceptualization, data curation, formal analysis, investigation, methodology, visualization, writing-original draft.

Dong-Wan Lee: PhD; contributed to data curation and investigation.

Su-Wan Lee: PhD candidate; contributed to data curation and investigation.

Nguyen Minh Thuyet: Professor, contributed to reviewing and editing the manuscript.

Nguyen Hoang Viet: Professor, contributed to reviewing and editing the manuscript.

Jin-Chun Kim: Professor, provided supervision, validation, funding acquisition, review, and editing of the final manuscript.

Acknowledgments

None.

Fig. 1.

Morphology of starting powders includes (a) gas-atomized Cu and (b) nano-Al₂O₃ powder, and (c) Cu-5 wt.% Al₂O₃ powder after mixing with EDS elemental mapping analysis.

Fig. 2.

Relative density of pure Cu and Cu-5 wt.% Al₂O₃ composite sintered at different temperatures.

Fig. 3.

OM micrographics of (a-c) pure Cu and (d-f) Cu-5 wt.% Al₂O₃ composite sintered at (a, d) 800 °C, (b, e) 900 °C, and (c, f) 1000 °C. The red and yellow arrows show pores and Al₂O₃ particles distribution in Cu matrix, respectively.

Fig. 4.

SEM micrographs (a-c) pure Cu and (d-f) Cu-5 wt.% Al₂O₃ composite sintered at (a, d) 800 °C, (b, e) 900 °C, and (c, f) 1000 °C.

Fig. 5.

EBSD maps of (a-c) pure Cu and (d-f) Cu-5 wt.% Al₂O₃composite sintered at 1000 °C: (a, d) image quality (IQ) maps; (b, e) grain boundary (GB) maps; (c, f) inverse pole figure (IPF) maps.

Fig. 6.

Grain size distribution of (a) pure Cu and (b) Cu-5 wt.% Al₂O₃ composite sintered at 1000 °C.

Fig. 7.

High-resolution EBSD analysis of Al₂O₃-rich region in Cu-5 wt.% Al₂O₃ composite sintered at 1000 °C: (a) IQ map; (b) phase map; (c) IPF map.

Fig. 8.

XRD patterns of pure Cu and Cu-5 wt.% Al₂O₃ composite sintered at different temperatures.

Fig. 9.

Vickers microhardness of pure Cu and Cu-5 wt.% Al₂O₃ composite sintered at different temperatures.

Fig. 10.

Apparent electrical conductivity of pure Cu and Cu-5 wt.% Al₂O₃ composite sintered at different temperatures.

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