Research on the Microstructure and Mechanical Properties of Cr₂O₃/Cu Composites Prepared by Internal Oxidation and HP Method

Abstract

In this study, an innovative internal oxidation-powder metallurgy combined process was employed to controllably generate nano-sized Cr₂O₃ reinforcing phases within the Cu matrix. The Cu/Cr₂O₃ composites were successfully fabricated using the hot-press sintering (HP) method, and a systematic comparison was made between the microstructure and mechanical properties of composites prepared by internal oxidation and external addition methods. The results show that internal oxidation primarily occurs during the sintering process rather than ball milling. Compared with external addition, the internal oxidation method effectively prevents particle aggregation and achieves a uniform distribution of Cr₂O₃ particles in the Cu matrix. When the Cr content reaches 5 wt%, the Cu-5%Cr composite exhibits optimal mechanical properties, with a yield strength of 282.7 MPa and ultimate tensile strength of 355 MPa, representing increases of 43% and 34% over pure copper, respectively, while maintaining an elongation of 12.6%. The Cr₂O₃ particles generated via internal oxidation enhance their strength through Orowan strengthening and dislocation pinning, thereby significantly improving mechanical performance without compromising plasticity. This research provides a novel process optimization approach for developing high-performance dispersion-strengthened copper matrix composites.

1. Introduction

Traditional copper alloy materials are increasingly unable to meet the diverse and sophisticated requirements of the modern industry for high-performance copper alloys, particularly in terms of mechanical properties, functional characteristics, and environmental adaptability [1,2,3]. To address these challenges, researchers have adopted the second-phase strengthening strategy to develop copper-based composite material systems. These systems typically consist of a continuous copper matrix reinforced with discrete phases. This design maintains the excellent electrical and thermal conductivity of the copper matrix while significantly enhancing the material’s hardness and wear resistance [4,5,6,7]. In recent years, oxide dispersion strengthening (ODS) technology has garnered significant attention due to its exceptional strengthening effects. By uniformly dispersing high-hardness oxide ceramic particles (such as Cr₂O₃, Al₂O₃, ZrO₂, TiO₂, etc.) within the copper matrix through mechanical alloying or in situ synthesis processes, multiple strengthening mechanisms can be synergistically employed to further improve material performance [8,9,10,11,12,13].

In the current research domain, the fabrication methods for ODS copper-based composites primarily include powder metallurgy, mechanical alloying, and internal oxidation [14]. Chromium oxide (Cr₂O₃), as a ceramic material with high hardness and superior wear resistance, has garnered significant attention in metal matrix composite applications [15,16]. Cr₂O₃ particles exhibit a corundum-type crystal structure (hexagonal system, R-3c space group), which not only imparts high hardness and thermal stability but also significantly enhances the mechanical properties of copper-based composites, making them an ideal choice for reinforcing particles [17]. Despite extensive research on Cr₂O₃ particle-reinforced copper-based composites, achieving the uniform distribution of Cr₂O₃ particles within the copper matrix and further optimizing the performance of these composites remain significant challenges in practical applications [18,19,20]. Currently, various fabrication processes are employed for copper-based composites, including mechanical alloying, internal oxidation, reactive spray deposition, powder metallurgy, and in situ reaction at the liquidus line [21]. However, traditional methods such as melting and powder metallurgy, while capable of integrating ceramic particles with the metal matrix, often struggle to precisely control the uniform distribution of particles, potentially leading to fluctuations in material performance [22,23]. The internal oxidation method, by introducing oxide particles within the metal matrix, effectively avoids particle agglomeration and achieves uniform particle distribution through the precise control of the oxidation reaction process [24]. Nevertheless, studies on the preparation of Cr₂O₃/Cu composites using the internal oxidation method are relatively limited, and controlling the oxidation reaction remains complex. The oxidation process requires the precise control of temperature, time, and atmosphere to ensure uniform oxide distribution. If not properly controlled, it may result in non-uniform particle distribution or agglomeration, adversely affecting the performance of the composites [25,26]. Therefore, in contrast to the inherent drawbacks of traditional mechanical alloying and powder metallurgy methods, which often result in non-uniform distribution of reinforcing phases, this research innovatively integrates the internal oxidation approach with hot-pressing sintering (HP) technology. This integration enables the in situ formation and sub-micron-scale homogeneous dispersion of Cr₂O₃ reinforcing phases within the copper matrix. As a result, it successfully overcomes the bottlenecks of interfacial contamination and particle segregation associated with the external addition method. Combining the internal oxidation method with HP for the fabrication of Cr₂O₃/Cu composites is expected to be a highly promising research direction for enhancing the performance of copper-based composites.

This study focuses on the preparation of Cr₂O₃/Cu composites using the internal oxidation method combined with HP, and investigates the impact of the dispersed distribution of Cr₂O₃ particles on the mechanical properties of copper-based composites. To comprehensively and systematically evaluate the performance of the composites, advanced characterization techniques such as X-ray diffraction (XRD), scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS), and transmission electron microscopy (TEM) are employed to conduct an in-depth analysis of the microstructure and mechanical properties [27,28]. Additionally, this study examines the effects of different Cr₂O₃ contents generated during the internal oxidation process on the material’s properties, thereby determining the optimal reinforcement ratio and providing a critical basis for material performance optimization. By comparing the performance of Cr₂O₃/Cu composites prepared via external addition and internal oxidation methods, this research evaluates the improvement effect of Cr₂O₃ particles fabricated by the internal oxidation method on the copper matrix properties, offering valuable theoretical references and practical guidance for the development of high-performance copper-based composites.

2. Experimental Methods

2.1. Raw Materials

The raw materials employed in this research encompass 5 μm pure copper powder (with a purity of 99.9%, supplied by Shanghai Naiou Nano Technology Co., Ltd., Shanghai, China), 5 μm chromium powder (with a purity of 99.99%, furnished by Shanghai Naiou Nano Technology Co., Ltd., Shanghai, China), and 4 μm Cu₂O powder (with a purity of 99.99%, provided by Shanghai Naiou Nano Technology Co., Ltd., Shanghai, China). In the mixing phase, pure copper powder and chromium powder were blended in the pre-determined ratios, where the mass fraction of chromium was 1%, 3%, 5%, and 7%, and the samples were, respectively, labeled as C1, C3, C5, and C7. The mass ratio of Cu₂O to Cr was set at 1.5:1 for the fabrication of the composite powder. The composition of the composite material is presented in

Table 1. The compositions of the raw powders and the theoretical values of Cr₂O₃ in the sintered composites.

Specimen Designation	Powders	Cr2O3 Composition in the Composites (wt%)
C1	Cu2O/Cu-1 wt% Cr	0.438
C3	Cu2O/Cu-3 wt% Cr	1.315
C5	Cu2O/Cu-5 wt% Cr	2.19
C7	Cu2O/Cu-7 wt% Cr	3.58
A5	Cu-2.19 wt% Cr2O3	2.19

.

2.2. Preparation of Materials

The fabrication process of the Cr₂O₃/Cu composite material, as illustrated in Figure 1, involves several key steps: powder blending, ball milling, pre-pressing, and internal oxidation sintering. In the mixing stage, powders are weighed and blended according to the proportions specified in Table 1. The mixed powders are then placed in a ball-milling jar with an appropriate amount of alcohol as a dispersant. Mechanical mixing and ball milling are performed using a PMQW planetary ball mill at a rotational speed of 300 rpm for 30 h, with a ball-to-powder mass ratio of 10:1. The prepared Cu-Cr metastable alloy powder was placed in an environment at 60 °C and dried in a drying oven for 8 h. Next, the dried composite powder is loaded into a tube furnace and heated at a rate of 5 °C/min to 400 °C under a reducing atmosphere of argon–hydrogen mixture for 5 h. According to the mass ratio of Cu₂O to Cr (1.5:1), the Cu-Cr mixed powder is ball milled with 4 μm Cu₂O powder for 15 h at 300 rpm, maintaining a ball-to-powder mass ratio of 10:1. For example, after the Cu-1 wt% Cr composite powder was prepared by mixing 30 g of Cu powder and 0.3 g of Cr powder, subsequent treatment was carried out. Based on the mass ratio of Cu₂O to Cr being 1.5:1, the mass of Cu₂O was determined to be 0.45 g. Similar procedures were applied to other samples. The composite powder is then placed in a uniaxial pressing mold pre-coated with boron nitride release agent and pressed at a pressure of 30 MPa to form a preformed blank. Subsequently, the pre-pressed blank is transferred to a vacuum hot-pressing sintering furnace and heated from room temperature to 850 °C at a rate of 10 °C/min. The temperature is held at 850 °C for 2 h while applying a pressure of 50 MPa to ensure concurrent internal oxidation and sintering, promoting particle contact and densification. Similar fabrication processes are employed for preparing Cr₂O₃/Cu composites using both the external addition method and the internal oxidation method for comparative experiments. The amount of Cr₂O₃ generated is calculated based on the following reaction equation:

XCu + 3Cu₂O + 2Cr→(6 + X)Cu + Cr₂O₃

(1)

The fabricated bulk sample (denoted as A5, see Table 1) was subjected to a comparative study with the bulk sample prepared by the internal oxidation method. To ensure consistency with the internal oxidation product, we procured Cr₂O₃ powder with particle sizes and theoretical values matching those of the internal oxidation process. This Cr₂O₃ powder was added to the copper matrix using the external addition method. Subsequently, the Cu-5.0 wt% Cr alloy powder was mechanically mixed and ball-milled using a PMQW planetary ball mill for 45 h at a rotational speed of 300 rpm, with a ball-to-powder mass ratio of 10:1. The milling process utilized carbide balls with diameters of 10 mm and 5 mm, in a large-to-small ball ratio of 65:35. The prepared Cu-Cr₂O₃ alloy powder was dried in a drying oven at 60 °C for 8 h. The technological process was identical to that of the internal oxidation method, encompassing two major steps: ball milling and sintering.

2.3. Experimental Methods and Characterization

The phase analysis of the powder before internal oxidation and the bulk sample after internal oxidation was conducted using an XRD (model: XRD-7000, manufacturer: SHIMADZU, Kyoto, Japan). The X-ray source was Cu-Kα radiation, with operating conditions of 40 kV and 30 mA, and a scanning rate of 0.02°/min. The morphological characteristics of the powder and bulk samples, as well as the tensile fracture surfaces, were examined using a SEM (SEM, model: Talos, manufacturer: FEI, Eindhoven, The Netherlands) equipped with an EDS. The microstructure of the Cr₂O₃/Cu composite material, particularly the distribution of Cr₂O₃ particles within the composite, was investigated using a TEM (TEM, model: Tecnai G2 F20, Eindhoven, The Netherlands, operating at 200 kV). Elemental analysis was performed using an energy-dispersive X-ray spectrometer (EDS, brand: Bruker, Karlsruhe, Germany) attached to the TEM. For TEM sample preparation, disk-shaped samples with a diameter greater than 3 mm were cut from the HP sintered samples and ground to a thickness of approximately 50 μm. Subsequently, the disk samples were polished at 4.3 keV using a precision ion-milling system (model: Gatan 691, Pleasanton, CA, USA) to achieve electron transparency. Vickers hardness was measured at room temperature using a microhardness tester (model: HV-1000 (A)) with a load of 50 g and a dwell time of 10 s. Hardness tests were conducted at 10 distinct positions, and the average value was determined. The electrical conductivity at room temperature was measured using an eddy-current conductivity meter (model: Sigma 2008 A, Xiamen, China), with each sample measured 10 times and the average value calculated. The relative density of the sintered composite material was tested using a precision electronic balance (model: CP114, Shanghai, China). The specimens were cut into horseshoe-shaped tensile specimens. These specimens were polished to dimensions of 26 mm in length, 1.96 mm in width, and 1.5 mm in thickness. Subsequently, an electronic universal testing machine (Model: UTM4204, Jinan, China) was employed to evaluate the tensile properties of the composite materials. The loading speed was set at 0.15 mm/min during the testing process.

3. Results and Discussion

3.1. Microstructural Analysis

Figure 2 features SEM images of Cu-5% Cr powders at different ball-milling durations. As shown in Figure 2a, during the initial stage of ball milling, the powder surface is rough, with large and unevenly distributed particles. With an increase in ball-milling time to 20 h, the frequency and duration of collisions between the balls and powder particles increase, leading to increased brittleness and a higher propensity for fracture [29]. As the ball milling progresses, work hardening causes the powder to become increasingly brittle. The continuous collisions and impacts result in particle fragmentation, gradually reducing the particle size (as illustrated in Figure 2b) [30]. Throughout this process, dislocations and lattice distortions accumulate within the powder. As ball-milling time increases, the intensity of collisions and impacts significantly intensifies, transforming the powder into multi-layered flaky particles under vigorous impacts, accompanied by cold welding between particles [31]. Although the particles have been significantly refined, unlike the multi-layered and ductile powders observed in Figure 2c, the particle size is notably reduced, and the grain refinement process is closely associated with particle fracture and cold welding [29,32]. As ball-milling time further extends, the powder particles undergo repeated deformation, cold welding, work hardening, and fracture. At 45 h of ball milling, a balance between welding and fracture is achieved, with an increased welding rate, a tendency for the average particle size to decrease, smaller and smoother powder particles, and a more homogeneous distribution (as presented in Figure 2d). Additionally, Figure 2d,e show that C5 powder exhibits a more uniform particle distribution and less agglomeration compared to A5 powder. The A5 powder may not have undergone complete mechanical alloying, resulting in larger particles and some degree of aggregation. To verify whether the Cu and Cr powders are fully mechanically alloyed, XRD analysis was conducted on the powders at various ball-milling durations.

Figure 3a shows the XRD patterns of Cu-5 wt% Cr mixed powders after ball milling for 10 h, 20 h, 30 h, and 45 h. In the Cu-5 wt% Cr powder mixture, due to the low Cr content, the primary diffraction peaks originate from pure copper. As the ball-milling time increases, the intensity of the diffraction peaks gradually decreases, indicating a reduction in grain size and refinement of the crystal structure. This is attributed to the increased grain boundaries and grain refinement caused by the mechanical energy input during ball milling. The width of the diffraction peaks broadens, which results from both the reduced grain size and increased lattice strain. After 45 h of ball milling, the main diffraction peak angle slightly shifts to the left (i.e., 2θ decreases), indicating an increase in the Cu lattice parameter. This aligns with the dissolution process of Cr atoms, whose larger atomic radius causes lattice expansion. Between 20 h and 45 h, significant changes occur in the diffraction peaks, particularly near the (111) crystal plane, suggesting that Cr elements gradually dissolve into the Cu matrix and may start forming solid solutions. After 45 h of ball milling, the intensity of the diffraction peaks weakens and the peak width increases, indicating an increased proportion of Cu-Cr solid solutions in the alloy system. However, these changes may also be related to grain refinement, lattice strain, and other structural modifications, implying that the ball-milling energy is approaching saturation and the formation of solid solutions may still be in a non-equilibrium or metastable state. According to Bragg’s law, the main diffraction peak angle of Cu alloys is slightly smaller than that of pure Cu. Due to the small difference in atomic radius between Cu and Cr, the peak shift is minimal. Therefore, the 2θ angle of the (111) diffraction peak of the Cu-5 wt% Cr powder alloy after 45 h of ball milling is 43.21°, slightly lower than 43.45° for pure Cu. If a small amount of oxygen or Cr atoms (with a larger atomic radius) dissolve into the Cu lattice, the lattice parameter may increase significantly, causing a leftward shift in the diffraction peak [33]. Another possibility is that ball milling may introduce stacking faults in the Cu lattice, leading to effects opposite to those caused by the dissolution of larger atoms [34]. The experimental results further confirm the accumulation effect of mechanical energy during ball milling, indicating that extending the ball-milling time promotes the formation of Cu-Cr solid solutions. After 45 h of ball milling, the accumulated mechanical energy is sufficient to drive the generation of Cu-Cr solid solutions. This conclusion is consistent with the findings of Aguilar et al., highlighting the critical role of continuous mechanical energy accumulation in the formation of Cu-Cr solid solutions [35]. As the ball-milling time increases, oxidation may occur (e.g., of Cu or Cr), leading to the formation of minor oxides. Although oxidation can affect the lattice structure, no new distinct peaks (such as those characteristic of CuO or Cr₂O₃) are observed in the spectrum, indicating that oxidation is minimal and has not significantly impacted the main crystal structure. The experimental results verify that the leftward peak shift trend is consistent with theoretical analysis, confirming that Cr atom dissolution leads to lattice expansion, in agreement with Bragg’s law and the physical properties of Cu-Cr alloys.

Figure 3b shows the XRD patterns of the Cu-5%Cr composite material, comparing the diffraction characteristics of the powder state (C5 powder curve) and the sintered bulk state (C5), with key crystal-plane peaks labeled. After internal oxidation treatment, a significant shift in the diffraction angle of Cu is observed, indicating that Cr has reacted with O to form Cr₂O₃ particles, which precipitate from the Cu lattice and distribute at the grain boundaries. Additionally, the width of the Cu(111) diffraction peak slightly broadens before oxidation, while the Cu₂O diffraction peak appears after ball milling (prior to oxidation). Notably, after oxidation, the Cu(111) peak width significantly narrows, the Cu₂O peak disappears, and no Cr₂O₃ diffraction peak is detected. The experimental results demonstrate that the internal oxidation process occurs during the sintering stage rather than the ball-milling stage. During sintering, Cr precipitates from the Cu solid solution and reacts with Cu₂O to form fine Cr₂O₃ particles. Due to the small size and low content of these Cr₂O₃ particles, their diffraction peak is difficult to observe in the XRD pattern. However, by comparing magnified images of local regions before and after oxidation, the Cr₂O₃ (024) diffraction peak can be clearly identified in the XRD pattern of the bulk material post sintering. The appearance of this Cr₂O₃ peak confirms that Cr undergoes internal oxidation during sintering, forming the reinforcing-phase Cr₂O₃. This peak does not appear in the powder sample, further confirming that the internal oxidation reaction primarily occurs during the sintering process.

Figure 4 presents the XRD patterns of pure Cu and Cu-Cr composites (C1, C3, C5, C7) after internal oxidation. The primary diffraction peaks correspond to the characteristic peaks of Cu, while no distinct Cr peaks are observed due to the low Cr content. In Figure 4b, the diffraction angle of Cu in the C1, C3, C5, and C7 alloy powders shifts to the right after internal oxidation. This indicates that Cr in the Cu-Cr alloy reacts with O to form Cr₂O₃ particles, which precipitate from the Cu lattice. The lattice constant of Cu decreases to 0.3615 nm, leading to an increase in the diffraction angle [36]. The lattice distortion caused by Cr solid solution in the Cu lattice is minimal, resulting in a subtle shift in the Cu peak after internal oxidation. The internal oxidation behavior can be confirmed by the shift in the Cu diffraction angle, consistent with the results reported in other literature [37]. In Figure 4a, the XRD pattern of pure Cu primarily exhibits characteristic diffraction peaks at the (111), (200), and (220) crystal planes. These peaks are intense and have a narrow FWHM, indicating a well-ordered crystal structure with large grain size and no presence of other phases. For the Cu-Cr composites (C1, C3, C5, C7), the characteristic Cu peaks at (111), (200), and (220) are retained. As the Cr content increases, the intensity of the Cu diffraction peaks gradually decreases, and the FWHM slightly broadens, suggesting a reduction in grain size, likely due to lattice distortion and particle refinement induced by ball milling. In composites with Cr content of C3 and higher, the characteristic peak of the Cr₂O₃ (024) crystal plane appears, indicating the formation of Cr₂O₃ particles within the composites. The intensity of the Cr₂O₃ peak progressively increases from C3 to C7, indicating that higher Cr content leads to more Cr₂O₃ particles.

Figure 5 presents the SEM images of powders with varying Cr contents and the theoretical Cr₂O₃ content generated by internal oxidation of 5% Cr, compared to powders with externally added Cr₂O₃ (Figure 5e). As the Cr content increases, the powder particles become finer, and the particle size distribution becomes more uniform. However, the dispersion of the externally added Cr₂O₃ powder is suboptimal due to poor wettability between Cr₂O₃ and the Cu matrix, which hinders effective bonding with Cu particles. Notably, the sample with 5 wt% Cr content (Figure 5c) exhibits significantly smaller particle sizes compared to other samples, indicating superior dispersion. Further analysis reveals that when the Cr content deviates from 5 wt% (such as 1 wt%, 3 wt%, and 7 wt%), the particle size increases (Figure 2d), suggesting that the external addition of Cr₂O₃ promotes particle agglomeration, reduces dispersion, and adversely affects the dispersion of Cu particles.

Figure 5a–d presents the SEM images and EDS mappings of Cu₂O/Cu-Cr composites with varying Cr contents. It can be discerned from Figure 5a–d that the Cr and O elements are homogeneously distributed within the copper matrix, signifying that the Cr₂O₃ particles produced via internal oxidation are uniformly dispersed in the Cu₂O/Cu-Cr composites (as depicted in Figure 5a–d). As the Cr content escalates from 1 wt% to 7 wt%, the black regions gradually augment and exhibit a denser distribution, indicating that the increase in Cr content exerts an influence on the microstructure of the samples [38]. Specifically, in the SEM images of the C5 (Cu₂O/Cu-5wt% Cr) and C7 (Cu₂O/Cu-7wt% Cr) composites (as illustrated in Figure 5c,d), distinct black areas emerge, whereas the Cu₂O/Cu-Cr composites with a Cr content lower than 5 wt% (as shown in Figure 5a,b) do not display evident black areas. To further investigate the composition of the black areas, in conjunction with the elemental analysis outcomes of the black areas in the Cu₂O/Cu-Cr composites, it is discovered that these black areas are primarily composed of aggregated chromium oxide particles. In the C5 and C7 composites, the chromium oxide particles are uniformly distributed. The variation in Cr content gives rise to significant disparities in the microstructure of the Cu₂O/Cu-Cr composites, and this microstructural alteration might impact the overall properties of the composites. Particularly in the presence of externally added Cr₂O₃, the black areas within the copper matrix significantly enlarge and present obvious particle aggregation, which might be associated with material-phase transformation, uneven element distribution, or precipitated particles.

Compared with C5 and A5, in Figure 5c, Cr is relatively dispersed. In Figure 5e, the distribution of Cr is more extensive and dense, indicating the aggregation of chromium oxide particles. The black spots observed by scanning electron microscopy (SEM) are chromium-containing oxide microparticles with a diameter of approximately 1 μm. To further explore the distribution of oxide particles within the composite material, an in-depth study of the microstructure of the Cu-Cr₂O₃ composite material was conducted using transmission electron microscopy (TEM). The elemental composition of the composite material was qualitatively analyzed by an energy-dispersive X-ray spectrometer (EDS, Bruker) equipped on the TEM, as shown in Figure 6.

Figure 6a,b presents the bright-field and dark-field transmission electron microscopy (TEM) images of the sintered Cr₂O₃/Cu samples, along with elemental analysis. Figure 6a shows the uniform dispersion of Cr₂O₃ particles in the Cu matrix, with uniformly refined particle sizes and clear grain boundaries, indicating a regular distribution of grains and grain boundaries. Figure 6c demonstrates the uniform distribution of Cr elements in the Cu matrix, indicating that the distribution of Cr in the sample is relatively uniform, confirming the successful dispersion of Cr₂O₃ particles during the preparation process. Additionally, through the magnification of the selected area electron diffraction (SAED) in the region of Figure 6a,d,e were obtained. Figure 6d displays the SAED pattern from the boxed region in (a); further confirming the existence of the Cr₂O₃ crystal structure. The clear diffraction spots indicate that the material has a good crystal structure without obvious crystal defects or amorphization. Figure 6e is a TEM image of the sample at a higher magnification, showing the macroscopic morphology of the large particles. From the figure, it can be seen that the Cr₂O₃ particles are nearly spherical, with a smooth surface and a hexagonal close-packed (HCP) crystal structure [39]. The average size of the Cr₂O₃ particles in Figure 6a,b is approximately 50 nm. HRTEM observations were carried out on the interface bonding area in Figure 6e, uncovering a tight interface bonding between the Cr₂O₃ particles and the Cu matrix, indicating a robust interfacial bonding force between the two phases. Figure 6f is a high-resolution TEM (HRTEM) image, depicting the microstructure of the crystal. From the figure, distinct lattice fringes can be clearly observed, with the marked lattice spacings of 0.4007 nm and 0.239 nm corresponding, respectively, to the interplanar spacings of Cr₂O₃ and Cu. This indicates that the Cr₂O₃ particles have been successfully embedded in the Cu matrix with satisfactory interface bonding. The clarity of the lattice fringes implies a high crystal quality of the material, without significant crystal defects.

The interplanar spacing of the Cu matrix is 0.239 nm, corresponding to the Cu (111) plane, while that of the Cr₂O₃ nanoparticles amounts to 0.4007 nm, corresponding to the Cr₂O₃ (012) plane. The mismatch parameter δ can be computed by the following formula: δ = (aα − aβ)/aα, where aα and aβ represent the lattice constants of the two phases, respectively. The calculated δ value is −0.677, suggesting a tensile mismatch between Cu (111) and Cr₂O₃ (012) particles. The greater the δ mismatch parameter is, the more severe the mismatch at the interface between the two phases becomes, the higher the interfacial energy is, and the stronger the inhibition of dislocation movement at this interface is. Considering the considerable mismatch at this interface, it implies that this interface can effectively impede dislocation slip. With the increase in the second-phase Cr₂O₃, the Cr₂O₃ particles are capable of hindering dislocation slip, thereby enhancing the yield strength and tensile strength of the material [33].

The clarity of the lattice fringes indicates a superior crystal quality of the material, with no significant crystal defects. Figure 6 comprehensively presents the microstructural characteristics of the Cr₂O₃/Cu composite material, ranging from particle distribution to crystal structure. The TEM and SEM images distinctly reveal the uniformity of the material, crystal quality, and interface bonding quality. The EDS analysis and HRTEM further verify the uniform distribution and embedded state of the Cr₂O₃ particles in the matrix. These results provide robust experimental support for the optimization of the preparation process and performance research of the Cr₂O₃/Cu composite materials.

3.2. Mechanical and Electrical Properties of Cr₂O₃/Cu Composites

Figure 7 illustrates the hardness, electrical conductivity, and density of Cr₂O₃/Cu composites.

Table 2. Hardness, electrical conductivity, and density of C1, C3, C5, and C7 composites.

Specimen Designation	Relative Density (%)	HV (N/mm2)	Conductivity (%IACS)
Cu	99 ± 0.1	91.15 ± 0.52	91.12 ± 0.43
C1	97.9 ± 0.2	118.25 ± 2.06	80.13 ± 0.52
C3	97.6 ± 0.3	129.25 ± 12.39	70.02 ± 0.29
C5	97.5 ± 0.3	137.25 ± 6.84	49.10 ± 0.46
C7	97.4 ± 0.2	181.25 ± 8.99	31.85 ± 1.14

provides detailed data on the density, electrical conductivity, and hardness of pure copper and sintered specimens C1, C3, C5, and C7 with varying Cr contents. As shown in the data, all samples exhibit densities exceeding 97% after sintering, indicating successful densification through HP sintering. With increasing Cr content, the internal oxidation process generates more Cr₂O₃ particles, which significantly enhances dispersion strengthening. Consequently, the hardness and strength of the sintered samples are markedly improved compared to pure Cu, highlighting the dispersion strengthening effect of Cr₂O₃ particles. Specifically, the hardness increases from 91.15 N/mm² for C1 to 137.25 HV for C7, demonstrating that higher Cr content leads to a greater amount of reinforcing-phase formation. The Cr₂O₃ particles generated during internal oxidation gradually refine and disperse uniformly, effectively hindering dislocation movement and enhancing material hardness and strength [33]. Although the mechanical properties improve, electrical conductivity decreases significantly from 91.1% IACS for C1 to 31.85% IACS for C7. This reduction in electrical conductivity can be attributed to the increased number of grain boundaries and the higher concentration of Cr₂O₃ particles generated by internal oxidation as Cr content rises. These factors enhance mechanical properties but impede electron mobility, thereby reducing electrical conductivity.

Figure 8a presents the tensile stress–strain curves of pure Cu at room temperature, the Cr₂O₃/Cu composites fabricated by the internal oxidation method, and the Cr₂O₃/Cu composites prepared by the external addition method as shown in Figure 8b. The results demonstrate that the ultimate tensile strength of pure copper is 268 MPa, with an elongation of 42.4%. The tensile strength of the Cr₂O₃/Cu composites is conspicuously higher than that of pure Cu. With the increase in the formation of Cr₂O₃, the tensile strength of pure copper rises from 268 MPa to 355 MPa for C5, but the elongation declines. This strengthening effect is primarily ascribed to the precipitation behavior of Cr elements during the internal oxidation process. This is because of the incoordinate deformation between Cr₂O₃ particles and the copper matrix, resulting in an enhancement of the tensile strength of the composites while sacrificing the elongation [40]. It can be known from Figure 6 that the TEM observation clearly reveals that Cr elements form fine Cr₂O₃ particles in the Cu matrix. With the increase in the formation of Cr₂O₃ particles, the tensile strength of the material is significantly enhanced, mainly for the following two reasons: (1) The dispersed Cr₂O₃ particles pin the dislocations. According to the dislocation theory, Cr₂O₃ particles possess good hardness and high-temperature stability. During the deformation process, when dislocations encounter Cr₂O₃ particles, they will bend and be obstructed. As the external stress increases, dislocation loops will form around the Cr₂O₃ particles. Under the stress, the dislocation loops become new dislocation sources, generating more dislocations and increasing the dislocation density, thereby improving the strength and hardness of the material. Additionally, when dislocation lines bypass the dislocation loops, they need to overcome the resistance of the particles and the reverse stress of the dislocation loops, increasing the work performed by the external force and enhancing the resistance to dislocation movement, thereby strengthening the tensile strength of the composite material [41]. (2) The Cr₂O₃ particles hinder grain-boundary migration, refine grains, and enhance the mechanical properties of the composite material. Specifically, they slow down the grain-boundary migration rate and increase the matrix strength. The finer and more uniformly dispersed the particles are, the greater the resistance to grain-boundary migration becomes. Nanoscale Cr₂O₃ particles are uniformly distributed within the copper matrix. During sintering, these hard Cr₂O₃ particles inhibit grain growth in the copper matrix, leading to finer grains. Although the tensile strength of the composite material has been improved, the influence of different preparation methods on its performance still requires further investigation. Figure 8b shows the tensile stress–strain curves of pure Cu, A5, and C5 samples. Pure copper exhibits typical ductile metal characteristics, with a fracture strain exceeding 40%, demonstrating excellent plastic deformation capacity. In contrast, the A5 composite material, prepared by the external addition method, shows a significant increase in tensile strength to approximately 220 MPa, but the fracture strain drops below 10%, indicating a marked decrease in plasticity. This reduction in plasticity may be attributed to uneven particle distribution or poor interface bonding. The C5 composite material, fabricated by the internal oxidation method, achieves a higher tensile strength approaching 350 MPa, while maintaining better plasticity compared to A5 but lower than pure copper. This suggests that the internal oxidation method provides a superior balance between strength and plasticity. In the A5 composite material, the Cr₂O₃ particles prepared by the external addition method may exhibit agglomeration or insufficient bonding with the Cu matrix, resulting in limited strengthening effects and increased brittleness. These issues could be related to uneven particle distribution during ball milling or sintering. In contrast, the C5 material prepared by the internal oxidation method features smaller and more uniformly distributed Cr₂O₃ particles, which effectively impede dislocation movement and significantly enhance strength. Additionally, the tighter interface bonding between Cr₂O₃ particles and the Cu matrix improves the plasticity of the C5 material compared to A5. Overall, Figure 8b clearly illustrates the balance between strength and plasticity. Pure copper demonstrates the best ductility but the lowest strength; the A5 material significantly increases tensile strength through the addition of Cr₂O₃ particles, but at the expense of plasticity; the C5 material achieves an optimal balance between strength and plasticity due to more uniform particle distribution, smaller particle size, and tighter interface bonding. A comparative analysis of Figure 8a,b reveals that the C5 material prepared by the internal oxidation method achieves a better balance between strength and plasticity, showcasing outstanding performance as a high-strength structural material. Compared to the external addition method, the Cr₂O₃ particles prepared by the internal oxidation method have a more uniform distribution, smaller particle size, and tighter interface bonding, which not only significantly enhances material strength but also retains ductility to some extent. Therefore, the internal oxidation method demonstrates superior comprehensive performance. Figure 8b further confirms that the strength of the C5 material is higher than that of A5, and its plasticity performance is better than that of A5. Although the A5 material prepared by the external addition method has elevated strength, its plasticity has decreased significantly due to uneven particle distribution and interface bonding issues.

The strain hardening rate of a material is defined as the derivative of true stress with respect to true strain. This parameter provides insight into the material’s ability to resist external loads during plastic deformation. When the strain hardening rate equals the true stress, the elongation is referred to as the uniform elongation. As shown in Figure 9a, the addition of Cr to Cu results in a significant increase in yield strength, while the elongation decreases. Figure 9b illustrates the strain hardening rate curves for both materials. To further analyze the variations in strain-hardening behavior, Figure 9b is divided at a true strain of 0.02, as depicted in Figure 9c,d. Figure 9c presents the initial deformation behavior of Cu and C5. It can be observed that after yielding, the strain-hardening rate of C5 initially increases before rapidly decreasing. In contrast, the strain-hardening rate of Cu remains relatively stable after yielding, followed by a gradual decline. The post-yield increase in the strain hardening rate of C5 contributes significantly to its enhanced yield strength. Additionally, the strain-hardening rate of C5 consistently exceeds that of Cu, indicating superior strain-hardening capability. This phenomenon can be attributed to Cr precipitates at the grain boundaries of Cu, which pin dislocations and hinder their movement. Consequently, the presence of Cr restricts plastic deformation in Cu, leading to precipitation strengthening and an increase in the overall strength and work-hardening capacity of C5.

As shown in Figure 9d, despite the higher initial work-hardening rate of C5, it exhibits a rapid decline in this rate during later stages of deformation, resulting in necking and reduced elongation. In contrast, Cu, unaffected by dispersed Cr, demonstrates a lower but more consistent work-hardening rate, which gradually decreases over time, thereby maintaining higher elongation.

Figure 10 presents a comparative analysis of the tensile fracture morphologies of pure Cu and two Cr₂O₃/Cu composites, C5 and A5, prepared using different fabrication processes. The examination of fracture features provides insights into the influence mechanisms of various preparation methods on the mechanical properties of the materials. In Figure 10a, the fracture surface of pure copper exhibits characteristic ductile fracture morphology, characterized by a large number of deep and uniformly distributed dimples. This microstructure indicates that pure copper undergoes substantial plastic deformation before fracture, leading to the formation of typical microvoid coalescence features [42]. Figure 10b reveals that the fracture surface exhibits obvious dimples and partial cracks, indicating a certain degree of ductility. The presence of Cr₂O₃ particles in the fracture zone suggests their involvement in crack initiation and propagation. The interfaces between the particles and the matrix may serve as crack initiation sites. Compared to pure copper, the C5 composite material fabricated via internal oxidation demonstrates an optimized balance between strength and plasticity. While the size of the dimples has decreased, evidence of plastic deformation is still observable. This indicates that the internal oxidation method enhances material strength while partially retaining its plasticity. Figure 10c shows that the fracture surface of the A5 composite material prepared by the external addition method is predominantly characterized by particle agglomeration and crack concentration regions, indicative of typical brittle fracture behavior. SEM observations reveal significant agglomeration of Cr₂O₃ particles. The agglomerated distribution of particles leads to stress concentration, accelerating crack initiation and propagation, ultimately resulting in material fracture. The reduction in dimple features suggests a weakened plastic deformation capability. Combined with the EDS analysis in Figure 5b, it is evident that there is pronounced elemental segregation in the particle agglomeration areas, confirming that the external addition method results in uneven particle distribution and poor interfacial bonding [43]. In contrast, the C5 material, fabricated via the internal oxidation method, exhibits a more uniform distribution of Cr₂O₃ particles. The fracture surface reveals an optimized balance between strength and plasticity. Conversely, the A5 material, prepared using the external addition method, shows a significant reduction in ductility due to pronounced particle agglomeration and poor interfacial bonding [41]. Through comparison, it is discovered that the primary advantage of the internal oxidation method resides in achieving the nanoscale dispersion of Cr₂O₃ particles and favorable interface bonding. This microstructure characteristic not only enables the reinforcing phase to effectively bear stress, effectively enhancing the strength of the matrix, but also retains a certain degree of plasticity. Nevertheless, as the external addition method relies on mechanical mixing, which has limitations and cannot completely avoid defects such as particle agglomeration and inadequate interface bonding; as such, the strength and toughness of the composite material are thus reduced. The refinement and uniform distribution of dimples further indicate that the internal oxidation method significantly improves the interface bonding performance between the particles and the matrix, thereby enhancing the fracture toughness of the material. This provides a solid theoretical basis for optimizing the preparation process of metal matrix composites.

4. Conclusions

In this study, Cr₂O₃/Cu composite were successfully fabricated using the internal oxidation method combined with HP. These composites were compared with those prepared by the external addition method. Through detailed observations of the microstructure and comprehensive performance evaluations, the following conclusions were drawn:

• Compared with the external particle addition method, the Cr₂O₃ particles synthesized in situ via the internal oxidation process exhibit uniform distribution. Additionally, the interfacial bonding strength between the Cr₂O₃ particles and the copper matrix is significantly enhanced, resulting in a superior balance of strength and plasticity for the Cr₂O₃/Cu composite. This also offers the possibility to construct multi-modal composite systems (such as Cr₂O₃-Al₂O₃ heterogeneous particles and graphene/Cr₂O₃ multi-dimensional structures) and optimizing the synergy between strength and toughness via interface engineering.
• The testing results of mechanical properties indicate that the Cr₂O₃/Cu composite material exhibits a markedly enhanced strength in comparison to pure Cu. As the quantity of Cr₂O₃ particles increases, the yield strengths of specimens C1, C3, C5, and C7 have increased significantly. Notably, specimen C5 shows a particularly remarkable improvement. Its yield strength has increased significantly to 282.7 MPa, and the ultimate tensile strength has reached 355 MPa. Compared with pure copper, these values represent increases of 43% and 34%, respectively. Simultaneously, this composite material demonstrates certain electrical conductivity properties, thereby showcasing the potential for the integration of structure and function.