The Microstructures and Properties of Cu-Ni-Co-Si Alloys: A Critical Review

Abstract

This review provides an overview of recent advancements in Cu-Ni-Co-Si alloys, focusing on their processing methods, microstructures, and properties. Due to their non-toxic composition, enhanced mechanical properties, and excellent electrical conductivity, Cu-Ni-Co-Si alloys have emerged as a promising alternative to traditional Cu-Be alloys in the electrical and electronics industry. This review discusses various synthesis techniques, including casting, vacuum induction melting, and additive manufacturing, and evaluates their effects on the formed microstructures. In addition, it explores the influence of different elements and thermal treatments on the alloys’ microstructures and properties, discussing strategies to enhance the properties of Cu-Ni-Co-Si alloys. Key strengthening mechanisms—including precipitation hardening, grain boundary strengthening, and solid solution hardening—are examined in detail, with particular emphasis on their synergistic effects in optimizing alloy performance. Furthermore, future research directions are highlighted, focusing on the optimization of alloying element concentrations and heat treatment protocols to achieve an enhanced balance between strength and electrical conductivity. These improvements are critical for meeting the demanding requirements of advanced applications in electronics and high-reliability components.

1. Introduction

Copper-based alloys have been one of the primary metallic materials utilized by humanity since ancient times. In recent years, copper alloys have gained even broader application across various industries, including electronics, automotive, and aerospace, due to their excellent mechanical, thermal, and electrical properties [1,2,3,4]. Among them, Cu-Be alloys have been the most widely used ultra-high-strength alloys in the past due to their superior mechanical and electrical properties [5,6]. However, Beryllium evaporation during melting poses significant environmental and health risks [7]. To address the toxicity issues associated with Cu-Be alloys, researchers have sought to develop new alloy materials, which led to the creation of copper-based alloys containing various trace elements. Examples include Cu-Ni-Al [8], Cu-Ti [9,10], Cu-Ni-Si [11,12], Cu-Cr [13], and Cu-Fe-P alloys [14].

In recent years, as electronic devices have evolved toward miniaturization and large-scale integration, the electrical industry has placed increasingly stringent demands on material performance for high-end products such as connectors, lead frames, and elastic components [15]. Previously mentioned copper alloys can no longer fully meet these performance requirements, spurring interest in Cu-Ni-Co-Si alloys due to their superior mechanical and electrical properties, as shown in Figure 1.

The emergence of Cu-Ni-Co-Si alloys marks a significant milestone in the development of copper-based alloys. Unlike Cu-Be alloys, Cu-Ni-Co-Si alloys are free of toxic elements, making them more environmentally friendly and easier to process. Studies have shown that after age hardening, this alloy surpasses Cu-Be alloys in both mechanical properties and electrical conductivity [16]. Additionally, Cu-Ni-Co-Si alloys exhibit excellent high-temperature softening resistance (higher than 500 °C), wear resistance, and corrosion resistance [17], broadening their application scope, particularly in the manufacturing of electronic components with high reliability requirements [17]. Today, Cu-Ni-Co-Si alloys are widely used in elastic components, connectors, and integrated circuit lead frames [18,19].

Figure 1. The comparison of properties between Cu-Ni-Co-Si [20,21,22,23], Cu-Ni-Si [24,25], Cu-Cr-Zr [26], Cu-Ti [27,28], Cu-Be [5,29], and Cu-Ni-Sn [30,31].

Figure 1. The comparison of properties between Cu-Ni-Co-Si [20,21,22,23], Cu-Ni-Si [24,25], Cu-Cr-Zr [26], Cu-Ti [27,28], Cu-Be [5,29], and Cu-Ni-Sn [30,31].

The primary advantage of Cu-Ni-Co-Si alloys lies in their combination of high strength and high electrical conductivity, which enables them to perform reliably in high-temperature environments (higher than 500 °C), making them suitable for electrical conductors and contact materials. In electronics, they outperform alternatives like Cu-Ni-Si by balancing strength and conductivity via nanoscale δ-(Ni,Co)₂Si precipitates, while their high-temperature stability (>500 °C) makes them ideal for aerospace components. Additive manufacturing leverages rapid solidification to achieve refined microstructures and near-net-shape fabrication, bypassing traditional processing limitations. Additionally, their sustainability advantages—such as reduced energy consumption and compliance with green regulations—position them as key materials for electric vehicles and renewable energy systems. Fundamentally, these alloys serve as a model for multi-mechanistic strengthening (precipitation + grain boundary + twinning), offering insights for advanced material design. However, these alloys also have certain drawbacks. For instance, enhancing the alloy strength often results in a reduction in electrical conductivity, a trade-off that limits their application in scenarios demanding extremely high conductivity [32,33,34]. Additionally, the machinability of these alloys is limited, especially at high strain rates (higher than 0.001 s⁻¹). Under these conditions, dynamic stress–strain curves show serrated oscillations, reflecting alternating softening and hardening [33,35].

The development prospects for Cu-Ni-Co-Si alloys are highly promising. Future research directions are expected to focus on optimizing the addition of microalloying elements and heat treatment processes to further enhance the alloy’s properties. Additionally, through the application of phase diagram calculations and kinetic models, a deeper understanding of the alloy’s precipitation processes and microstructural evolution can be achieved. This will enable precise control over the alloy’s performance, making it a pivotal area of future research [33,36,37]. Furthermore, the development of more efficient and energy-saving manufacturing processes to meet industrial demands for both performance and cost-effectiveness is another key research direction.

Building on a review of the existing literature on Cu-Ni-Co-Si alloys, this paper first introduces the advancements in Cu-Ni-Co-Si alloys over the past decade. It provides a brief overview of the alloy’s preparation processes and offers a detailed discussion of its microstructure, as well as the strengthening mechanisms underlying its superior mechanical and electrical properties. The influence of different trace elements is also examined. Finally, the paper explores in detail how various cold-working and heat treatment processes impact the mechanical and electrical performance of the alloy.

2. Diverse Manufacturing Approaches of Cu-Ni-Co-Si Alloys

The manufacturing of Cu-Ni-Co-Si alloys can be accomplished using various methods, including vacuum melting, forging, and additive manufacturing. This section will introduce each of these manufacturing methods in detail.

2.1. Casting

Casting is a conventional processing method for Cu-Ni-Co-Si alloys, where horizontal continuous casting and down-draw semi-continuous casting are commonly used [38]. While many high-strength, high-conductivity copper-based alloys for industrial applications have predominantly been produced by forging in recent years due to their cost-effectiveness and high throughput, casting remains an essential manufacturing route. However, to achieve the high strength and high conductivity required for these alloys, post-casting processes such as hot working, cold working, and heat treatment are often necessary. For applications such as lead frame materials, where the alloy needs to be fabricated into thin sections, large-scale variable processing techniques are typically employed for both cold and hot deformation.

Liu et al. [39] applied hot rolling to improve the properties of the alloy after casting and observed that under high-temperature deformation conditions, the alloy exhibited remarkable dynamic recrystallization behavior. This dynamic recrystallization led to grain refinement and enhanced the mechanical properties of the alloy. The hot rolling process is generally conducted within a temperature range of 760 °C to 970 °C, and by adjusting the deformation rate, the grain structure can be optimized to improve both the strength and ductility of the alloy. Zhu et al. [40] used an as-cast Cu-Ni-Co-Si alloy in the temperature range of 700–950 °C and the strain rate range of 0.01–5 s⁻¹ to finish the Isothermal hot compression experiment, founding that the softening of the Cu-Ni-Co-Si alloy during hot compression was influenced by discontinuous dynamic recrystallization (DDRX) and twinning mechanisms. The occurrence of DDRX was promoted by the increasing in Σ3 grain boundaries.

Following hot rolling, researchers often employ multi-step aging and multi-step cold rolling processes to further enhance the alloy’s properties. For instance, Huang et al. [22] utilized a two-step cold rolling and two-step aging process (Pre-cold-rolled by 70% reduction and aged at 400 °C for 2 h, then second cold rolled by 70% reduction and final aged at 400 °C, 450 °C, 500 °C for 5, 20, 30, 60, 120, 240 and 300 min), respectively, which resulted in a tensile strength of 1086 MPa and an electrical conductivity of 30.1% IACS. After undergoing two cold rolling stages, the alloy’s cross-sectional structure was no longer composed of fully recrystallized equiaxed grains, but instead exhibited elongated, highly deformed grains. This processing route significantly enhanced the conductivity with only a minor reduction in strength, and the presence of numerous annealed and deformation twins within the alloy contributed to the overall improvement in its performance.

2.2. Additive Manufacturing

Additive manufacturing (AM) is a technique that enables the rapid fabrication of complex geometric structures through the layer-by-layer fusion of metal powders. This method holds significant potential for improving material properties. In the Laser Powder Bed Fusion (LPBF) process, the rapid cooling of the melt pool leads to grain refinement, resulting in smaller grain sizes compared to traditional casting methods, and a more uniform distribution of precipitates. This, in turn, greatly enhances the strength of the alloy [41,42]. Moreover, LPBF allows for precise control over the alloy’s microstructure by adjusting parameters such as laser power, scanning speed, and powder deposition thickness [43]. Bartosz et al. [43] used LPBF to produce the alloy (laser power equal to 300 W, scanning speed of 500 mm/s, and powder deposition thickness of 0.12 mm), and the relative densities were increased to over 98% by optimizing the printing parameters, which resulted in a relatively coarse, columnar crystalline structure. But when the energy density is increased (laser power more than 500 W), the densities can reach more than 99%, and the grains become refined isometric crystals. Furthermore, subsequent aging and heat treatment transformed grains into equiaxed crystals with a peak yield strength of approximately 590 MPa and electrical conductivity of 34.2% IACS. For instance, during the additive manufacturing process, the precipitation phases of Ni and Si form finer Ni₂Si and (Ni,Co)₂Si phases under high cooling rates, which improves strength without sacrificing electrical conductivity [44].

One of the key advantages of additive manufacturing is its ability to offer flexibility in adjusting the alloy composition during the manufacturing process. By altering the proportions of alloy powders, localized strengthening or the enhancement of specific material properties can be achieved. For example, in the production of Cu-Ni-Co-Si alloys, varying the content of Ni or Co allows for the adjustment of the balance between electrical conductivity and mechanical properties. Additionally, additive manufacturing supports the design of gradient materials by gradually adjusting the composition layer by layer, thus imparting different performance characteristics to various regions of the alloy and enabling the production of multifunctional materials [45].

However, the printed alloy surface is rougher and exhibits inferior properties compared to cast alloys, so a subsequent heat treatment process (like solution and homogenization) is needed to improve the surface roughness and enhance the alloy’s properties. To improve the surface finish of Cu-Ni-Co-Si alloys, post-processing treatments such as laser remelting and heat treatment are often required. These processes help to enhance surface smoothness and overall material performance [46].

2.3. Vacuum Melting

Vacuum melting is primarily used for the production of high-performance alloys with high purity and low oxygen content (lower than 0.005%) [47,48,49]. This method minimizes the presence of impurities such as oxygen, nitrogen, and hydrogen in the melting environment, thus preventing the formation of excessive oxides and nitrides in the alloy. Similarly to additive manufacturing, vacuum melting allows for precise control over grain size and the distribution of precipitates by adjusting the temperature gradients during the melting and cooling processes. Slower cooling rates promote grain growth and improve the material’s ductility, whereas rapid cooling refines the grains, enhancing the material’s hardness and strength [50].

Additionally, vacuum melting supports multiple re-melting cycles, where the alloy undergoes several melting and re-solidification steps. This process further purifies the material and optimizes its microstructure, ensuring better uniformity. Especially in large-scale production, this approach guarantees consistent performance across different batches of material [50,51,52].

3. Microstructure of Cu-Ni-Co-Si Alloys

3.1. Precipitation Behavior

Lei et al. [53] investigated the precipitation behavior of Cu-Ni-Si alloys and described the evolution of precipitate phases under different heat treatment temperatures and durations. They also plotted the TTT (Time–Temperature–Transformation) curve for the Cu-Ni-Si alloy. As shown in Figure 2, the study revealed that at 400 °C, no precipitates formed in the alloy. Instead, the alloy underwent spinodal decomposition, leading to the formation of Ni- and Si-enriched clusters. As the heat treatment time increased, these clusters gradually grew into precipitates of β-Ni₃Si and β-NiAl phases, with particle sizes around 10 nanometers. With further extension of the aging time up to 8 h, these β-Ni₃Si and β-NiAl precipitates continued to grow and eventually reaching a size of several hundred nanometers.

For heat treatment in the range of 450 °C to 500 °C, the authors proposed that the precipitation sequence involved the oversaturated solid solution undergoing L1₂ ordering to form the β-Ni₃Si phase. Early in the aging process, Ni₃Si was the predominant phase. As the aging continued, the δ-Ni₂Si phase appeared, and eventually, both phases coexisted in the matrix. At 550 °C, early-stage aging led to the formation of discontinuous, banded precipitates (δ-Ni₂Si and δ-Co₂Si), primarily located at grain boundaries, with fine β-Ni₃Si particles distributed within these bands. At temperatures above 700 °C, precipitation was mainly dominated by continuous δ-Ni₂Si phases, which rapidly grew with time, increasing from several nanometers in the early stages to several hundred nanometers as aging progressed.

Figure 2. Phase transformation map of the Cu-Ni-Si alloy annealed at temperatures of 400–750 °C [53].

Figure 2. Phase transformation map of the Cu-Ni-Si alloy annealed at temperatures of 400–750 °C [53].

The primary precipitate phase in Cu-Ni-Co-Si alloys is δ-(Ni, Co)₂Si, which typically appears as nanometer-sized rod-like or disk-like structures in the matrix, with sizes generally ranging from 5 to 10 nanometers [54], as shown in Figure 3 The crystal structure of these precipitates is orthorhombic, and although there is some lattice mismatch with the Cu matrix, the mismatch-induced stress is relatively low [53], which contributes to enhancing the alloy’s strength while maintaining a degree of ductility (with tensile strength at about 800 Mpa, ductility at about 12%). The introduction of Co effectively promotes the formation of the (Ni, Co)₂Si phase by substituting a portion of the Ni atoms, leading to the formation of a more stable precipitate phase. During the aging process, as the aging time increases, the number and size of the precipitates gradually increase and become uniformly distributed within the matrix [22].

Figure 3. Detailed microstructure featuring precipitation δ-Ni₂Si [54].

Figure 3. Detailed microstructure featuring precipitation δ-Ni₂Si [54].

Huang et al. [22] verified this conclusion using transmission electron microscopy (TEM). As shown in Figure 4, the δ-(Ni, Co)₂Si phase is dispersed throughout the Cu matrix, with precipitates typically measuring tens of nanometers in size, and most of them exhibit rod-like or disk-like shapes. Huang et al. [55] focused on nano-scale precipitate coarsening during over-aging, observing that the core–shell structure consisted of a stable δ-Ni₂Si core and a metastable interface composed of Cu, Ni, and Si elements for the shell. But at the over-aged stage, the core–shell structure completely transformed into a stable and coarsened δ-Ni₂Si phase, remaining coherent with the Cu matrix [56]. With the point of diffusion kinetics, this kind of structure was caused by the disparity between the diffusion activation energy and the diffusion rate of Ni, Si, and Cu [57,58,59].

Figure 4. TEM images of the samples aged at 450 °C for 150 min and 500 min (Directly aged): (a) Convergent Beam Electron Diffraction image of the sample aged for 150 min viewed along the zone axis of [001]Cu; (b) HRTEM and IFFT images of precipitates in the sample aged for 150 min; (c) CDF image of the sample aged for 500 min viewed along the zone axis of [001]Cu; and (d) HRTEM and IFFT image of precipitate in the sample aged for 500 min [22].

Figure 4. TEM images of the samples aged at 450 °C for 150 min and 500 min (Directly aged): (a) Convergent Beam Electron Diffraction image of the sample aged for 150 min viewed along the zone axis of [001]Cu; (b) HRTEM and IFFT images of precipitates in the sample aged for 150 min; (c) CDF image of the sample aged for 500 min viewed along the zone axis of [001]Cu; and (d) HRTEM and IFFT image of precipitate in the sample aged for 500 min [22].

Figure 5. The microstructure and TEM images of the alloys after peak aging: (a,a1) alloys after 30% cold rolling; (b,b1) alloys after 50% cold rolling; (c,c1) alloys after 70% cold rolling; (d,d1) alloys after 80% cold rolling [60].

Figure 5. The microstructure and TEM images of the alloys after peak aging: (a,a1) alloys after 30% cold rolling; (b,b1) alloys after 50% cold rolling; (c,c1) alloys after 70% cold rolling; (d,d1) alloys after 80% cold rolling [60].

3.2. Dislocation Structure

As shown in Figure 5, Cu-Ni-Co-Si alloys—due to the multiple heat treatment and cold working processes they undergo—develop complex dislocation structures such as dislocation tangles, dislocation walls, and dislocation cells. The matrix exhibits dislocation tangling and precipitates surrounding dislocations resulting from cold rolling, as depicted in Figure 5(a1–c1). Upon peak aging of the CR-80 sample (Figure 5d), the increased aging temperature causes fast sub-grain boundary movement, leading to extensive recrystallized grain formation. Yet, a few nanometer-sized subgrains persist, as shown in Figure 5(d1).

Dislocation entanglements and dislocation lines are commonly observed after cold deformation, particularly when the alloy undergoes cold rolling at low temperatures, such as in liquid nitrogen rolling. These dislocation structures contribute to the strengthening of the material [61].

After certain annealing and aging treatments, dislocation cell structures have been observed. These dislocation cells form under high-strain-rate conditions through dislocation slip and climb mechanisms, which further enhance the hardness of the alloy. The presence of these dislocation substructures plays a significant role in improving the material’s mechanical properties, particularly its strength and resistance to plastic deformation.

Many previous studies have achieved excellent mechanical and electrical properties via regulating the thermomechanical treatment to control the precipitates and dislocation densities [61]. They found that with the increase in the aging temperature, the dislocation density presents a gradual decrease. Qi et al. [62] produce different dislocation densities in different aged samples by cold rolling, they try to determine a critical dislocation density to boost the nucleation sites of precipitates, to simultaneously maximize the mechanical and electrical properties of alloys.

3.3. Twinning

Wei et al. [61] observed that, in addition to dislocation cells and dislocation walls, a significant amount of twinning structures, including annealing twins and deformation twins, exist within Cu-Ni-Co-Si alloys. As shown in Figure 6, through high-resolution transmission electron microscopy (HRTEM) analyses, the researchers found that twins predominantly appear in alloys subjected to low-temperature cyclic rolling and subsequent aging treatments. Specifically, in samples aged at 350 °C for 2 h, a widespread layered twin structure was observed within the grains, closely associated with grain boundaries. These twin interfaces not only act as barriers to dislocation slip but also enhance the strength of the alloy through additional twin/slip mechanisms.

The formation of twins is closely related to the alloy’s low stacking fault energy (SFE). Studies have shown that the low SFE in Cu-Ni-Co-Si alloys promotes twinning [63,64]. During the aging process, as the temperature increases, the stacking fault energy gradually decreases, which enhances the formation of twins. The essence of a twin crystal is a region of the crystal where the atomic faces are stacked in mirror-symmetric order, and the stacking laminar error is a deviation from the local stacking order. When the layer dislocation energy is low, the energy barrier for the formation of layer dislocations decreases, and the low SFE makes the energy cost of the twin boundaries relatively lower. Reduced layer faults can make Shockley partial dislocations more susceptible to extension, leading to the generation of multiple stacking layer faults (multi-layer faults). From the crystallographic definition, single-layer stacking order anomalies (like ABCBCA) are called intrinsic layer faults, whereas ≥3-layer periodic mis-stacking (like ABCBACBA) constitutes a twin boundary. Thus, the increased width of extended dislocations in low-layer error-energy materials makes the critical twinning thickness (typically 3–5 atomic layers) more attainable, thus enhancing the twinning tendency.

Youssef et al. [63] discussed how the low SFE of the Cu alloy greatly affects its mechanical properties and deformation mechanism such that the ultimate stress of nanocrystalline Cu alloy is higher than nanocrystalline Cu. This mechanism plays a crucial role in explaining the relationship between the probability of twin formation and the aging temperature in these alloys. Twins, therefore, contribute significantly to the strengthening of the material by impeding dislocation motion and facilitating additional strengthening mechanisms.

4. Effect of Elements on Properties and Microstructure of Alloys

Recent advances in Cu-Ni-Co-Si alloy research have revealed critical insights into microstructural control and thermomechanical response. Particularly, the introduction of trace alloying elements has been demonstrated to simultaneously improve both mechanical strength and electrical conductivity—a traditionally challenging trade-off in copper alloys.

Table 1. Properties of Cu-Ni-Co-Si alloys with different Ni, Co, and Si ratios.

Composition/wt.%				Aging Process		Tested Properties
Cu	Ni	Co	Si	Temperature/°C	Time/h	MH/HV	EC/% IACS
Bal.	0.32	0.59	0.62	450	1	291	33.4	[65]
Bal.	0.32	0.61	1.09	450	1	262	42	[54]
Bal.	2.31	1.18	0.82	500	1	280	41	[66]
Bal.	0.85	2.67	0.82	500	1	250	46	[66]
Bal.	0.17	3.36	0.84	500	1	240	47.5	[55]
Bal.	1.16	2.30	0.82	500	1	258	44.5	[55]
Bal.	2.5	0.5	0.7	450	1	230	44.8	[67]
Bal.	2.5	0.5	0.7	450	2	220	46.7	[56]
Bal.	2.3	0.7	0.7	450	1	232	31.7	[56]
Bal.	2.3	0.7	0.7	450	2	248	38.3	[56]
Bal.	2.0	1.0	0.7	450	1	267	43.9	[56]
Bal.	2.0	1.0	0.7	450	2	248	44.6	[56]
Bal.	2.5	0.5	1.0	450	4	248	37.0	[50]
Bal.	2.5	0.5	1.0	450	3	249	38.0	[39]
Bal.	2.3	1.2	0.8	450	1	275	40.2	[39]
Bal.	4.5	1.0	1.0	400	2	329	30	[39]
Bal.	1.5	1.0	1.5	450	1	200	43	[68]
Bal.	2.0	1.5	0.9	450	3	233	37.3	[57]
Bal.	1.0	2.5	1.8	450	2	233	52.5	[69]
Bal.	8	0.5	1.2	450	2	268	30.3	[58]
Bal.	7.0	0.1	1.6	450	2	316	15.9	[58]
Bal.	4.5	1.2	1.0	400	1	263	24.9	[22]
Bal.	4.5	1.2	1.0	450	0.5	280	26.7	[17]
Bal.	4.5	1.2	1.0	450	1	329	30.1	[17]
Bal.	63.2	0	25.76	450	4	273	30	[70]
Bal.	51.43	0	27.55	450	4	248	16	[59]
Bal.	1.82	0.62	0.86	450	1	265	42.5	[71]
Bal.	1.82	0.62	0.86	450	2	258	44	[60]
Bal.	1.28	0.84	0.46	450	1	300	35	[72]
Bal.	1.5	0.7	0.7	450	2	260	40.5	[73]

Table 1. Properties of Cu-Ni-Co-Si alloys with different Ni, Co, and Si ratios.

Composition/wt.%				Aging Process		Tested Properties
Cu	Ni	Co	Si	Temperature/°C	Time/h	MH/HV	EC/% IACS
Bal.	0.32	0.59	0.62	450	1	291	33.4	[65]
Bal.	0.32	0.61	1.09	450	1	262	42	[54]
Bal.	2.31	1.18	0.82	500	1	280	41	[66]
Bal.	0.85	2.67	0.82	500	1	250	46	[66]
Bal.	0.17	3.36	0.84	500	1	240	47.5	[55]
Bal.	1.16	2.30	0.82	500	1	258	44.5	[55]
Bal.	2.5	0.5	0.7	450	1	230	44.8	[67]
Bal.	2.5	0.5	0.7	450	2	220	46.7	[56]
Bal.	2.3	0.7	0.7	450	1	232	31.7	[56]
Bal.	2.3	0.7	0.7	450	2	248	38.3	[56]
Bal.	2.0	1.0	0.7	450	1	267	43.9	[56]
Bal.	2.0	1.0	0.7	450	2	248	44.6	[56]
Bal.	2.5	0.5	1.0	450	4	248	37.0	[50]
Bal.	2.5	0.5	1.0	450	3	249	38.0	[39]
Bal.	2.3	1.2	0.8	450	1	275	40.2	[39]
Bal.	4.5	1.0	1.0	400	2	329	30	[39]
Bal.	1.5	1.0	1.5	450	1	200	43	[68]
Bal.	2.0	1.5	0.9	450	3	233	37.3	[57]
Bal.	1.0	2.5	1.8	450	2	233	52.5	[69]
Bal.	8	0.5	1.2	450	2	268	30.3	[58]
Bal.	7.0	0.1	1.6	450	2	316	15.9	[58]
Bal.	4.5	1.2	1.0	400	1	263	24.9	[22]
Bal.	4.5	1.2	1.0	450	0.5	280	26.7	[17]
Bal.	4.5	1.2	1.0	450	1	329	30.1	[17]
Bal.	63.2	0	25.76	450	4	273	30	[70]
Bal.	51.43	0	27.55	450	4	248	16	[59]
Bal.	1.82	0.62	0.86	450	1	265	42.5	[71]
Bal.	1.82	0.62	0.86	450	2	258	44	[60]
Bal.	1.28	0.84	0.46	450	1	300	35	[72]
Bal.	1.5	0.7	0.7	450	2	260	40.5	[73]

Research has shown that within a specific Ni/Co ratio range (usually at 3.5~4.5), as shown in Table 1, the mechanical performance and conductivity of the alloy can be synergistically optimized [66]. This section outlines the trace elements added to Cu-Ni-Co-Si alloys and describes their effects on the alloy’s microstructure and performance.

4.1. Ni and Si

Cu and Ni exhibit complete solid solubility, with Ni’s atomic radius being ~3% smaller than Cu’s. Therefore, when Ni atoms enter the Cu matrix lattice, the size difference induces lattice distortion, which hinders dislocation motion and improves the alloy’s strength [54,74]. However, according to Matthiessen’s rule [75], when a significant number of Ni atoms dissolve into the Cu matrix, the electrical resistivity of the alloy increases sharply. As a result, the Ni content is generally controlled within 3%.

In contrast, Cu and Si are not completely soluble, so the amount of Si added is typically less than that of Ni, usually controlled below 1.5%. The contribution of Si to alloy strength is similar to that of Ni. When Si atoms enter the Cu matrix, they also induce lattice distortion, making dislocation slip more difficult and ultimately improving the alloy’s strength. Additionally, the solid solution strengthening effect of Si interacts synergistically with that of Ni [76], further enhancing mechanical properties. Ni and Si also contribute to precipitation strengthening, as Ni and Si combine to form fine precipitates, such as Ni₂Si [77]. These precipitates hinder dislocation motion through the Orowan mechanism, further enhancing the alloy’s mechanical strength. This effect is particularly pronounced after aging treatment, where the precipitates are uniformly distributed in the matrix, significantly hardening the alloy [78]. The non-deformable nanoscale precipitated phase enhances the strength of the alloy by impeding the movement of dislocations, which are forced to bypass these particles in slip, bending, and reconnecting while leaving dislocation rings around the particles when the precipitated phase is uniformly distributed after the potentization process [79,80]. This process requires additional stress, and subsequent dislocations have to overcome greater resistance (as the dislocation rings increase the effective particle size), which ultimately results in the hardening of the alloy. The strengthening effect is dependent on particle spacing (smaller is stronger) and size (nanometer scale is optimal) and is the primary means of strengthening age-hardening alloys.

Similarly to Ni, the presence of Si creates barriers in the matrix, impeding the free movement of electrons and reducing electrical conductivity. Si atoms reduce electrical conductivity primarily through electron scattering [81,82,83]. As solute atoms, Si distorts the Cu lattice, increasing defect scattering (Matthiessen’s rule). When precipitated as SiO₂ or silicide, phase boundaries further scatter electrons (Mayadas–Shatzkes model). The effect intensifies with Si content; even 0.1 wt.% Si significantly reduces conductivity. In age-hardening alloys, optimized aging can partially restore conductivity by reducing solute Si, but it remains below pure Cu’s level due to residual precipitates. At the same time, some studies have shown that Si in Cu-Ni-Co-Si alloys not only contributes to precipitation strengthening but also refines the grains [76,84,85,86], promoting a uniform grain distribution. Refined grains help improve the yield strength and creep resistance of the alloy, particularly under high-temperature conditions, where the refined microstructure plays a crucial role in maintaining alloy stability.

4.2. Co

Adding trace amounts of Co to Cu-Ni-Si alloys has been a major breakthrough in the field of copper alloys for electronic components in recent years. The introduction of Co, combined with further heat treatment, can introduce fine precipitates of β-Ni₃Si and δ-Ni₂Si phases, which are typically coarse in Cu-Ni-Si alloys, and gradually transition their morphology to disk-like or fine rod-like shapes [53]. The addition of Co modulates the lattice mismatch between the matrix and the precipitated phase, and the (Ni,Co)₂Si phase in the Cu-Ni-Co-Si alloys has a lower mismatch with the Cu matrix than δ-Ni₂Si, which reduces the strain energy, thus suppressing the coarsening of the precipitated phase and sustaining a smaller size. Furthermore, Elemental Co hinders the diffusion of Ni and Si, and the bias of Co in the matrix reduces the diffusion rate of solute atoms, retards the Ostwald ripening process (precipitated phase growth mechanism), and results in a more homogeneous and finer distribution of precipitated phase sizes.

Additionally, since both δ-Ni₂Si and δ-Co₂Si belong to the orthorhombic crystal system, Co and Ni can co-form nanoscale δ-(Ni, Co)₂Si phases. These precipitates aggregate and distribute uniformly in the matrix. Jiang et al. [23] confirmed this through three-dimensional atom probe (3DAP) results (as shown in Figure 7). At the nanoscale, δ-Ni₂Si and δ-Co₂Si phases are observed to co-aggregate, which can be understood as a complex composite phase. These precipitates preferentially nucleate and grow at grain boundaries, and their morphology helps stabilize the structure, reducing interface energy and thereby minimizing the system’s total free energy, resulting in a more stable alloy with enhanced mechanical properties [87,88]. Xiao et al. [35] found that the electrical conductivities of the Cu-1.4Ni-1.2Co-0.9Si samples are generally 5–15% IACS higher than those of the Cu-2.8Ni-0.6Si samples processed by the same heat treatment. For example, the electrical conductivities of the Cu-2.8Ni-0.6Si samples aged at 450 °C and 550 °C for 4h are 37% and 43% IACS, respectively, while the Cu-1.4Ni-1.2Co-0.6Si samples are 49% and 47% IACS, respectively. As a result, Cu-Ni-Co-Si alloys demonstrate exceptional performance in applications requiring both high strength and high conductivity, such as in the electronics industry and modern integrated circuits.

Adding Co not only promotes the formation of more precipitates, thereby improving the mechanical properties of the alloy, but also inhibits the spinodal decomposition of δ-Ni₂Si and β-Ni₃Si in the early stages of aging, which reduces the hindrance to electron mobility within the matrix and enhances the electrical conductivity of the alloy [70,89,90]. Li et al. [91] employed first-principles calculations based on a density functional theory study to understand how trace elements contribute to alloys’ strength and electrical conductivity. The result shows that doping Co atoms into the copper matrix exhibited the most significant improvement in both mechanical and electrical properties, with the Young’s modulus, hardness, and electrical conductivity all achieving a huge boost compared to pure cooper Furthermore, the adding of Co effectively reduced electron losses at the interface between the precipitation and the matrix. The electron transmission rate through the (211)_Co2Si/(111)_Cu interface reached 78.6%, representing a 94.1% improvement compared to the (111)_Ni2Si/(111)_Cu interface.

Figure 7. Three-dimensional atom probe (3DAP) maps (5 × 5 × 15 nm) of solute atoms in the precipitates aged at 500 °C for different times: (a,e) 5 min, (b,f) 15 min, (c,g) 1 h, and (d,h) 4 h [23].

Figure 7. Three-dimensional atom probe (3DAP) maps (5 × 5 × 15 nm) of solute atoms in the precipitates aged at 500 °C for different times: (a,e) 5 min, (b,f) 15 min, (c,g) 1 h, and (d,h) 4 h [23].

4.3. Ce (Cerium)

The rare earth element Ce forms fine Ce-rich phases within the alloy, which can effectively increase the density of dislocations and inhibit grain boundary migration [92,93]. These Ce-rich phases exhibit excellent thermal stability in high-temperature environments, and through the particle-stimulated nucleation (PSN) mechanism, they delay the recrystallization process at elevated temperatures.

In the Cu-Ce system, the particle-excited nucleation (PSN) mechanism relies mainly on the local strain concentration generated during thermo-mechanical processing of micrometer-sized CeCu₅/CeCu₂ intermetallic (≥0.8 μm), which promotes recrystallisation through the formation of geometrically necessary dislocations (GNDs) density gradients and subgranular evolution (critical strain ε ≥ 0.6) Grain formation, which is most pronounced at 400–600 °C and low strain rates (<0.1 s⁻¹), is inhibited by nanoscale Ce phases (<200 nm) through the Zener pinning effect [94,95]. The lower PSN efficiency of Cu-Ce compared to the Al-Ce system is mainly attributed to the higher interfacial energy and the stronger tendency of the Cu substrate to dynamically revert [96]. Meanwhile, the in situ characterization of synchrotron radiation reveals that the formation of the >15° orientationally differentiated grain boundaries in the antipodal figure of merit (IPF) of the PSN region is directly related to the formation of the (111)_Cu/(0001)_CeCu5 orientation relationship. The core issues that remain unresolved in this field include the accurate determination of the CeCu₅/Cu interfacial energy (DFT deviates from the experimental value by up to 15%) and the role of trace Si additions in the modulation of the PSN mechanism [97].

This leads to a finer chromium particle size and more uniformly distributed recrystallized grains, which are also fine. This thereby improves the alloy’s high-temperature strength and thermal stability [98]. The interaction of dislocations with chromium precipitation phases is the main factor contributing to the excellent mechanical properties of alloys with the addition of Ce [99,100].

4.4. Mg, Cr and Zr

The amounts of Mg, Cr, and Zr added to Cu-Ni-Co-Si alloys are typically very small (less than 0.1%). Due to their low concentrations, the influence of these elements on alloy performance is not as significant as the effects of Ni, Co, Si, or Ce. These elements are soluble in the Cu matrix, similarly to Ni, and can form nanoscale precipitates with other elements, such as Si. These precipitates effectively suppress grain growth and grain boundary migration, thereby increasing the alloy’s recrystallization temperature and improving its high-temperature performance [101,102].

Mg plays a role analogous to that of Ce in metallic systems. It effectively pins dislocations, thereby promoting dynamic recrystallization and leading to the formation of finer, more uniformly distributed grains. The incorporation of Mg not only enhances the alloy’s strength and ductility but also significantly improves its post-cold-working tensile properties [103,104].

A previous study found that as the Cr content increased, the stress relaxation resistance of alloys improved [72,101,105]. Adding 0.15% Cr can produce Cr precipitates with a non-coherent orientation relationship with a Cu matrix, which are effective in pinning dislocations. Furthermore, the Cr element also increased the grain size, which reduced the preferential orientation of the grains and the activation of the dislocation sources during microplastic deformation [106].

Figure 8. Comparison of properties of alloys with different Zr contents after aging for different times at 400 °C, 450 °C, and 500 °C: (a–c) electrical conductivity (d–f) hardness [107].

Figure 8. Comparison of properties of alloys with different Zr contents after aging for different times at 400 °C, 450 °C, and 500 °C: (a–c) electrical conductivity (d–f) hardness [107].

The elements Cr and Zr are often added to copper alloys simultaneously. Since Zr has an extremely low solubility in Cu and co-precipitates with Cr during aging, this significantly refines the size of the precipitated phases [73,108]. Furthermore, Zr tends to segregate at the interfaces of precipitated phases, inhibiting their coarsening and thereby improving the alloy’s resistance to over-aging. As shown in Figure 8, after aging at 450 °C for 6 h, the peak hardness of Cu-3.2Ni-0.7Si-0.05Zr reaches 270.1 HV, with the conductivity of 37% IACS. Cu-Ni-Co-Si alloys usually reach peak hardness with a 60 min aging process, which is much lower than Cu-Ni-Si-Zr alloys. Moreover, the addition of Cr and Zr greatly increases the recrystallization temperature of copper alloys, further improving their resistance to softening [106,107].

5. Strengthening Mechanisms

5.1. Grain Refinement Strengthening

In addition to the regulation of trace elements, heat treatment processes play a critical role in enhancing the performance of Cu-Ni-Co-Si alloys. Studies have observed the crystalline phases of the alloy in its as-cast state and after homogenization, hot rolling, and solution treatment. Compared to the as-cast state, which exhibits a large number of dendritic structures, and the fragmented grains formed after hot rolling, solution treatment results in refined equiaxed grains. These grains contain a high density of annealing twins (more than 30%), contributing significantly to the strengthening of the alloy.

Furthermore, the use of thermomechanical deformation techniques, such as hot compression, can improve the recrystallization behavior of the alloy, leading to further grain refinement. As discussed earlier, the addition of various trace elements can delay the recrystallization process, suppress grain growth at high temperatures, and promote grain refinement.

According to the Hall–Petch relationship, refined grains significantly enhance the strength of the alloy by increasing the number of grain boundaries, which act as barriers to dislocation motion. This mechanism ensures that grain refinement remains a crucial factor in improving both the strength and thermal stability of Cu-Ni-Co-Si alloys.

5.2. Solid Solution Strengthening

The enhancement in yield strength and tensile strength due to the formation of solid solutions between elements such as Ni, Si, and the Cu matrix arises from lattice distortion caused by differences in atomic sizes. This strengthening effect can be quantified using the following equation [52]:

$$∆{\sigma }_S=\sum G\delta x_a^{2/3}\sqrt{{\displaystyle \frac{x_a}{3}}}$$

where $\delta$ is a factor of lattice change, $x_a$ is the atom fraction of the element, and $G$ is the shear modulus.

The dissolution of solute atoms into the matrix inevitably affects the properties of the alloy. However, in Cu-Ni-Co-Si alloys, this is generally undesirable because an increased number of impurity atoms in the matrix significantly reduces the alloy’s electrical conductivity. However, solid solution strengthening can effectively enhance the mechanical properties of alloys. However, the associated reduction in conductivity may be unsuitable for certain applications such as integrated circuits, lead frames, and Chip packaging strips. The precise contribution of solid solution strengthening to the overall strengthening mechanism requires further investigation.

As a result, the extent to which solid solution strengthening contributes to the overall strengthening of the alloy remains a subject of further investigation. Balancing mechanical performance and electrical conductivity is a critical challenge, and the role of solid solution strengthening in this balance requires careful consideration.

5.3. Precipitation Strengthening

Cu-Ni-Co-Si alloys are typical precipitation-strengthened alloys, where the second-phase precipitates serve as the primary strengthening mechanism. Based on the Orowan mechanism, the strengthening effect can be expressed by the following equation [53,54]:

$$∆{\sigma }_{\mathrm{O}\mathrm{r}\mathrm{o}\mathrm{w}\mathrm{a}\mathrm{n}}={\displaystyle \frac{0.91MGBln\left(d_p/b\right)}{\pi {\left(1-v\right)}^{\frac{1}{2}}\left(d_p\sqrt{\frac{3\pi }{2f_v}}-d_p\right)}}$$

where $M$ is the Taylor factor, $G$ is the shear modulus of the matrix, $d_p$ is the average diameter of particles, $b$ is the Burgers vector, $v$ is the Poisson’s ratio, and $f_v$ is the volume fraction of particles.

Various studies have demonstrated through calculations that nanoscale, uniformly distributed precipitates make a significant contribution to the mechanical properties of alloys [109,110,111,112]. One of the objectives of different heat treatment and cold working processes is to promote the precipitation of such nanoscale precipitates, which not only enhances the alloy’s mechanical performance but also improves its electrical properties.

5.4. Dislocation Strengthening

Wei et al. [61], using X-ray diffraction, measured and statistically analyzed the dislocation density and stacking fault probability of twin structures in Cu-Ni-Co-Si alloy samples under different aging and cold rolling processes, as shown in Figure 9. Figure 9a shows the diffraction patterns of Cu–Ni–Si–Co alloy aged at 250 °C, 300 °C and 350 °C for 2h measured by HEXRD. The data points in the different diffraction peaks are extracted, as shown in Figure 9c, and linear fitting is performed to obtain the slope. It can be seen from the results of Figure 9b that the error of the twin stacking-fault probability determined by the modified WH method is relatively large.

Their research explored the relationship between dislocation density and aging temperature, revealing that dislocation structures have a notable impact on alloy performance. They also concluded that substructures such as twins and dislocations contribute to strengthening the alloy through mechanisms similar to those of dislocation strengthening.

Figure 9. (a) HEXRD diffraction signals of Cu–Ni–Si–Co alloy with 250 °C, 300 °C, and 350 °C, (b) The linear fitting of the modified Williamson-Hall method based on the experimental results of HEXRD, (c) The linear fitting of the modified Williamson-Hall method based on the experimental results of HEXRD [61].

Figure 9. (a) HEXRD diffraction signals of Cu–Ni–Si–Co alloy with 250 °C, 300 °C, and 350 °C, (b) The linear fitting of the modified Williamson-Hall method based on the experimental results of HEXRD, (c) The linear fitting of the modified Williamson-Hall method based on the experimental results of HEXRD [61].

The accumulation of dislocations increases the density of barriers to dislocation motion, enhancing the alloy’s yield strength and hardness. Combined with the formation of twin boundaries, these features create a synergistic effect, further improving the alloy’s mechanical properties.

6. Future Works

Future research on Cu-Ni-Co-Si alloys should prioritize optimizing alloy composition and processing techniques to further enhance their mechanical and electrical properties. Key areas include refining microalloying strategies (e.g., Cr, Zr, Ce) to stabilize nanoscale precipitates and suppress coarsening, as well as developing advanced thermomechanical processes like cryorolling or hybrid additive manufacturing to tailor microstructures. Forging and vacuum induction melting remain economical and widely used processes, providing refined grain structures and uniform properties. Additive manufacturing, although still in its exploratory stages, has demonstrated significant potential for enhancing alloy performance due to its precise control over microstructure and composition.

Future research should focus on a systematic analysis of these elements to further optimize alloy performance. Proper heat treatment methods are crucial to achieve the desired phase distribution, and missteps in thermal processing could negatively impact the overall performance of the alloy. Computational modeling and in situ characterization will be crucial for understanding interface dynamics and defect interactions, enabling precise control over performance.

Additionally, efforts should focus on improving high-temperature stability, fatigue resistance, and sustainability, ensuring these alloys meet the demands of next-generation electronics and aerospace applications. Scalability and cost-effectiveness must also be addressed to facilitate industrial adoption. By bridging fundamental research with practical engineering, Cu-Ni-Co-Si alloys can overcome current trade-offs and expand their role in high-performance technologies.

Future research directions should address several critical aspects: First, the fundamental relationship between crystallographic texture evolution during processing and its dual impact on both mechanical strength and conductivity pathways requires systematic investigation. Second, emerging computational methods, particularly machine learning algorithms, show significant potential for accelerating alloy development by predicting optimal composition–processing–property relationships. These approaches can overcome the traditional trial-and-error limitations in designing next-generation high-performance alloys.

7. Conclusions

In this section, a preliminary summary of the Cu–Ni–Co–Si alloy is presented in terms of diverse manufacturing approaches, strengthening mechanisms, and potential applications. This review systematically examined the interplay between microstructure and properties, focusing on processing methods, strengthening mechanisms, and the roles of key alloying elements (as shown in

Table 2. Summary of main findings in this review.

Summary of Major Findings
Aspect	Key Findings	References
Processing Methods	-Additive manufacturing (LPBF) achieves finer grains and uniform precipitates. -Vacuum melting reduces impurities, enhancing ductility. -Hot/cold rolling combined with aging optimizes the strength-conductivity balance.	[22,44,46,57]
Microstructure	-δ-(Ni, Co)2Si precipitates (5–10 nm) dominate strengthening. -Dislocation cells/twins form during thermomechanical processing. -Low stacking fault energy promotes twinning.	[9,35,36,40]
Alloying Effects	-Co: Increases conductivity by 5–15% IACS vs. Cu-Ni-Si. -Ce: Refines grains, inhibits recrystallization. -Cr/Zr: Raise recrystallization temperature, resist over-aging.	[76,99]
Strengthening Mechanisms	-Precipitation hardening contributes >70% of yield strength. -Grain refinement (Hall-Petch) and dislocation density are critical for strength. -Twin boundaries impede dislocation slip.	[23,44]
Future Directions	-Optimize heat treatment protocols (e.g., multi-step aging). -Explore machine learning for texture/property relationships.

).

The conclusions of this review are as follows:

(1) As a copper-based high-strength and high-conductivity alloy, Cu-Ni-Co-Si can be effectively utilized in advanced electronic components, connectors, springs, and integrated circuit lead frames. Regarding preparation processes, forging, vacuum induction melting, and additive manufacturing each contribute uniquely to the optimization of the alloy’s mechanical properties and electrical conductivity.

(2) The key factors influencing the properties of Cu-Ni-Co-Si alloys include grain refinement, precipitation of strengthening phases, and the distribution of the δ-(Ni, Co)₂Si phase. Microalloying elements such as Co, Si, and Ce play crucial roles in enhancing the mechanical properties and electrical conductivity through mechanisms like precipitation strengthening and grain boundary strengthening. However, there are research gaps regarding the effects of microalloying elements such as Zr, Mg, and Cr, particularly in understanding their influence on recrystallization and grain boundary behavior.

(3) The primary challenge in developing advanced Cu-Ni-Co-Si alloys lies in achieving the simultaneous enhancement of mechanical properties while maintaining electrical conductivity above 45% IACS (International Annealed Copper Standard). Current research strategies focus on three key approaches: (1) the precise addition of trace alloying elements (Co, Cr, and Zr) to optimize precipitation strengthening; (2) advanced thermomechanical processing techniques, including controlled hot/cold working to refine microstructure; and (3) the meticulous control of precipitate size distribution and volume fraction through tailored aging treatments.