We first investigate the mechanical behaviors of the Cu/Ni NMMCs with various tilt angles. Figure 2 shows the classical stress-strain curves of the corresponding loading for different tilt angles. It can be observed from Fig. 2 that for θ = 0° , the yield stress is highest. For θ > 0° , the yield stress decreases with increasing degree, until it reaches 33° and then it continuously increases. Meanwhile, the curve of 45° displays an increase in stress with increasing strain until the yield strain, followed by fluctuations around the critical stress, implying that the plastic response of 45° is much better than the other samples. It can be seen from Fig. 2 that there are finite initial stresses in the “ a” region, although the samples are relaxed adequately before tensile loading. This is mainly because the interfacial and surface stress cause an intrinsic compressive stress in the interior of the nanoscale metallic multilayer composites, and the compressive stress should be overcome by a stress at the initial moment of the tensile loading process. In addition, figure 2 shows that the Young’ s modulus of the Cu/Ni nanoscale metallic multilayer composites slightly decreases with increasing twist angle. It can be explained that for an fcc crystal, E(100) < E(110) < E(111), ^{[31, 32]} which means that the Young’ s modulus reaches a maximum value in 〈 111〉 and a minimum value in 〈 100〉 . For the sample with a relatively low tilt angle (for example 10° ), the crystalline orientation almost does not deviate from the (111) plane, while for high twist angle (for example 45° ), the crystalline orientation deviates from the (111) plane severely. Hence, the Young’ s modulus decreases with the twist angle increasing from 0° to 45° .

Fig. 2.

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	Fig. 2. Typical stress– strain curves for the Cu/Ni multilayer nanowires in different angles under tension loading.

To reveal the deformation mechanisms for the sample at various angles under uniaxial loading, the dynamic evolution of the dislocation is analyzed. Figure 3 shows the atomic configurations of Cu/Ni multilayer nanowires with 0° and 33° at different strains of loading. In the case of θ = 0° , the first dislocation nucleates from the ledge of the surface in the softer layer that is Cu in this case as shown in Fig. 3(a). Hence, the crystal structure of the Ni layer is still nearly perfect without lattice defect. After nucleation, the dislocations propagate toward the Cu layer interiors along 〈 111〉 . As the sample is further loaded to 0.12 strains, multiple slip systems are activated in the layer, and the dislocations intersect and cross over each other, thus enhancing the hardening effect, and strengthening the nanowires. In the case of θ = 33° , the dislocation emits from the surface of the Cu layer, and travels along a slip plane until it reaches the interface. With continuing to load, two slip systems are activated and dislocations with less different slip systems may intersect and cross with each other, thus leading to the softening behavior of material as seen in Fig. 3(b). The strength of material is controlled by dislocation interactions.

Fig. 3.

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	Fig. 3. Atomic configurations with different twist angles: (a) the Cu/Ni nanoscale metallic multilayer composites with a twist angle of 0° under different tension strains of 0.04 and 0.12; (b) Cu/Ni nanoscale metallic multilayer composites with a twist angle of 33° under different tension strains of 0.035 and 0.115. Atoms are colored by using a common neighbor analysis method (hcp in red).

Figure 2 shows that the model of 45° displays excellent plasticity response. For comparison, the atomic configurations of 45° and 10° tilt angle modes at the same strain levels are also shown in Figs. 4(a) and 4(b), respectively. Under 45° tension, most partial dislocations nucleate in the intersection region of the interface and surface in the initial stage, and glide on the interface planes, on the layer with the fewest defects as shown in Fig. 4(a). As the loading increases, the interface migration resulting from partial dislocation gliding on interface planes (see Fig. 4(a)) plays an important role in the plastic deformation.^{[21, 33]} In contrast, for the case of 10° , plenty of intersecting dislocations and interface planes are observed; a higher density of stored dislocations than that for the 45° case is shown in Fig. 4(b). This difference is caused by the crystalline orientation that dominates each of the slip systems.^[18] In the case of 45° , the favorable slip systems are coplanar with the interface, the slip directions are parallel to the interface and hence dislocations intersecting with interface planes are rarely seen here. Therefore, the slip system leads to the decrease of dislocation pileup near the interface, which reduces the stress concentration to a large extent, thus leading to the increase in plasticity of material. In the 10° case, the slip plane and the slip direction of the active slip system incline to interface, thereby inducing the dislocation– dislocation and dislocation– interface interaction. From Fig. 4(b), numerous dislocations pile up and interaction induces a large material potential energy, which hinders the plastic deformation of material.

Fig. 4.

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	Fig. 4. Atomic structure configurations of Cu/Ni with tilt angle θ = 45° (a) and θ = 10° (b) at the same strain levels.

In order to understand the underlying mechanism of the 45° , we also investigate the deformation behaviors of the structure with different layer thickness values. Figure 5 shows the stress– strain curves for NMMCs with layer thickness (h) values of about 1.88 nm, 3.76 nm, 5.01 nm, and 6.98 nm. From Fig. 5, we can find that the Cu/Ni NMMCs with 45° tilt angle displays excellent plasticity, regardless of layer thickness. This result indicates that the twist interface plays an important role in determining the plasticity properties of Cu/Ni nanowires. It can also be seen from Fig. 5 that the effect of layer thickness on the Young’ s modulus is very slight, the same as that in the 45° case. In addition, the simulation result indicates that the yield strength can be evidently enhanced with increasing layer thickness. Thus there may exist a reversal of the Hall– Petch (HP, the smaller grain size, the higher strength) effect. This is consistent with the conclusion obtained by Yuan and Wu^[25] in Cu/Ag multilayer nanowires. To further understand the deformation mechanism, the atomic structures of Cu/Ni nanowire with layer thickness values of 1.88 nm and 6.98 nm under the critical strain are shown in Figs. 6(a) and 6(b), respectively. For the nanowire sample with a larger thickness value, there will be more storage room for dislocations. As the storage room is too small for dislocations to pile-up inside the lattice, the mechanical properties will be very limited. It follows from Fig. 6 that the capacity of material for dislocation storage with a layer thickness of 1.88 nm is obviously lower than that with a layer thickness of 6.98 nm. Furthermore, the dislocation storage capacity of NMMC does not increase with increasing layer thickness when layer thickness is larger than 5.01 nm, and the material properties change slightly. The result is mainly due to the fact that the magnitude of the repulsive force and the room for dislocation storage can be affected by the change of layer thickness.^[1]

Fig. 5.

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	Fig. 5. Typical stress– strain curves for the samples of 45° with different layer thickness values under tension loading.

Fig. 6.

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	Fig. 6. Atomic structure configurations of Cu/Ni with layer thickness (a) h = 1.88 nm and (b) 6.98 nm under their corresponding critical stresses. Arrows show the interactions of dislocations in different slip planes.

Subsequently, we analyze the microscopic structure. Figure 6(a) shows partial dislocations nucleated from the intersections of the free surface and interface plane and moves in the direction parallel to the interface in the case of h = 1.88 nm at the yield point. Owing to the confinement imposed by the small size layers, there is limited room for the dislocations to move along the slip plane, leading to most dislocations being parallel to the interface, and rarely being seen to intersect with the interface plane. In contrast, when the layer thickness becomes 6.98 nm, figure 6(b) shows that more dislocations nucleate from the interface in the Cu layer, thereby activating the different slip systems, and dislocations interact more with each other, resulting in the highest yield stress. This result is in agreement with a previous atomistic result, ^[24] where it is shown that the slip systems improve the properties of the material.