Mechanical properties optimization and microstructure analysis of pure copper gradient laminates via rolling

doi:10.21203/rs.3.rs-1729350/v1

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Mechanical properties optimization and microstructure analysis of pure copper gradient laminates via rolling

https://doi.org/10.21203/rs.3.rs-1729350/v1

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Heterostructured (HS) material with extraordinary mechanical properties has been regarded as one of the most promising structural materials. In this paper, heterogeneous lamella structure (HLS) Cu laminates composed of surface elongated-grain layer and center equiaxed-grain layer were fabricated via rolling bonding and annealing. In order to study the interface effect on mechanical properties of gradient structured materials，both the laminate metal composite (LMC) and the non-composite laminate (NCL) were fabricated by cold-rolling pretreatment of the center layer (60% reduction) and cold-rolling bonding of the whole blank (67% reduction). Then by controlling the post-annealing regimes, the HLS was obtained and the microstructure of each layer was optimized, and a larger degree of microstructure heterogeneities such as grain size, misorientation angle and grain orientation were obtained, which resulted in obvious mechanical differences. Tensile tests of HLS, surface layer, center layer and NCL specimens reveal that the HLS annealed at 300°C/1h has a significantly higher strength than the center layer and a higher elongation than the surface layer. The HLS have a tensile strength and elongation at fracture of 278.08 MPa and 46.2%, respectively, which shows good balance of strength and plasticity. The main reason for the improved properties is the strengthening or strain hardening generated by the inhomogeneous deformation of the heterogeneous layers in the laminate and the mutual constraint acquired by the distinct layers with strong mechanical differences. The HLS has an interfacial bonding strength of 178.5 MPa, which plays a vital role in the coordinated deformation of heterogeneous layers. The method of obtaining gradients in homogeneous materials proposed in this study provides a new way of thinking and enriches the theoretical basis for the comprehensive performance improvement of traditional metallic materials.

Rolling

Heterogeneous lamella structure

Strength-ductility synergy

Interface

Copper is abundant in nature, stable in nature, and easy to mine and store. It has excellent properties such as better electrical conductivity, thermal conductivity, ductility, corrosion resistance and wear resistance, and is widely used in electric power, electronics, energy, machinery, metallurgy, transportation, light industry, new industries, and other fields. With the rapid progress of high-tech industries, new materials are being developed to improve their strength and plasticity. It becomes increasingly necessary to improve the performance of pure copper materials. As is the case with other traditional metals, pure copper is subject to the inverse law "strength increase-toughness decrease" [1,2]. It is this contradictory relationship between strength and toughness that severely limits the further enhancement of the overall performance of pure copper.

In recent years, heterostructured materials have been a research hotspot and extensive efforts are devoted to tailoring the microstructure of materials, including nano-twinned structure [3,4], bi-modal structure [5-7], gradient structure [8-12] and heterogeneous lamella structure [13-18], to achieve superior strength-ductility combination. Cheng et al. [4] fabricated gradient nano-twinned structures of Cu with a dual gradient of twin thickness and grain size by direct-current electrodeposition, achieving a combination of extremely high yield strength of 434 MPa and uniform elongation of 9.2%. They found through systematic experiments and atomistic simulations that this unusual behavior is afforded by a unique patterning of ultra-high densities of dislocations in the grain interiors. Yonghao Zhao et al. [7] prepared bimodal Cu with homogeneous distribution of different-volume-fraction micro-grains via isothermal recrystallization of an ultrafine grained Cu matrix at 200℃. Analysis of the tensile results suggests that the bi-modal Cu follows the rule of mixtures, with tensile strength of 273 MPa and elongation at fracture of 42% after 90 min of low temperature oil bath annealing. K.Lu [11] have used a surface mechanical grinding treatment (SMGT) for preparing a NG Cu film with a spatial gradient in grain size on a bulk CG Cu substrate and have achieved large tensile plasticity in the NG structure and revealed a different governing deformation mechanism. It is generally recognized that materials with nano-twin structures and gradient structures have exceptional mechanical properties, however, their preparation is both complicated and costly, and only materials of a specific size can be prepared [19]. Bi-modal material can be prepared by annealing-induced secondary recrystallization of UFG materials to improve the plasticity of UFG materials, but it is difficult to precisely control the nucleation sites and volume distribution of recrystallized grains during preparation [7]. Nano-twined, gradient and bi-modal materials are limited to large scale, high volume production applications due to the defects in their preparation.

The interactions and constraints between the different layers in a laminated material lead to co-deformation, and this allows the composite material to combine the excellent properties of the layers to achieve better mechanical properties [20]. Considering the preparation of laminated composites, various processing methods such as vacuum hot pressing (VHP) [21], explosive welding [22], high pressure torsion (HPT) [23] and rolling [24] have been proposed. For the VHP process, the good metallurgical combination by diffusion between layers can be realized under high temperature and pressure. But the cost of vacuum condition is high, and the size of products is limited by vacuum chamber. Explosive welding is a special method of pressure welding, which uses the detonation wave produced in the process of explosion to combine the homogeneous or heterogeneous metal materials. However, there are some issues in explosive welding, such as poor controllability and serious pollution. HPT is an attractive severe plastic deformation (SPD) method since it can refine the microstructure of laminated composites significantly, but the sample size processed by HTP is relatively small, which is not beneficial to practical industrial application. Compared with the above processes, as a kind of continuous processing technology, rolling has simple process conditions and can be used to produce large-scale plate structural materials. The technology of obtaining composite plates by rolling is widely used in industrial production.

For a long time, scientists at home and abroad have been seeking to develop high strength and ductility materials, and in recent years, great progress has been made in breaking the inverse law of "increasing strength and decreasing toughness" to acquire metals with well-matched strength and toughness [25-28]. By electrodeposition, surface mechanical grinding (SMGT), high pressure torsion (HPT), and equal channel angle extrusion (ECAP), the microstructure of the material can be altered to improve its strength without harming its toughness and plasticity. However, these methods are not suitable for high-volume industrial production. We hope to obtain laminates with heterogeneous laminate structures that can be industrially produced in large quantities through the rolling process to improve the comprehensive mechanical properties of metallic materials. In this paper, we constructed pure copper laminates with heterogeneous microstructures by rolling and annealing processes, and analyzed their mechanical properties through tensile and tensile-shear experiments, demonstrating their excellent mechanical properties and exploring the best preparation method.

2.1 Materials Preparation

The materials used in this experiment are T2 pure copper plates (Cu, 99.9 wt.%) with different thicknesses. The length × width × thickness of the two T2 pure copper plates are 100 mm × 30 mm × 2 mm (copper plate A) and 100 mm × 30 mm × 5 mm (copper plate B), respectively. The chemical composition is shown in Table 1.

Table 1

Chemical compositions of raw materials
Chemical compositions (wt.%)
Cu	Bi	As	Sb	Fe	Pb	S	Others
99.9	0.0008	0.0016	0.0017	0.00304	0.0038	0.0041	<0.1

2.2 Preparation processing

Figure 1 displays the process for preparing pure copper laminates with heterogeneous microstructure. Firstly, the copper plate B is cold-rolled at a reduction of 60% in order to obtain the initial blank copper plate C. Copper plate A and copper plate C, as the original blanks for the surface and central layers, maintained approximately equiaxed grains with an average grain size of 12.4 µm and elongated grains (length-width ratio 7.7:1) with an average grain width size of 5.6 µm, respectively. Then it is divided into two groups for billet assembly and rolling. In the first group, we will remove the oxidation layer, decontaminate, and sand the surfaces of two copper plates A and one copper plate C to be laminated, stack them in the order of copper plate A/copper plate C/copper plate A, so that the laminated surfaces of copper plate A and copper plate C are in complete contact, and bind their ends with aluminum wire. After binding, the 6mm thick original billet is cold-rolled through a two-roller mill at 67% reduction to obtain a 2mm thick laminate metal composite (LMC) with a strong bond. In the second group, we decontaminated the surfaces of two copper plates A and one copper plate C to be laminated, and fully coated the surfaces with isolating agent, stacked cold-rolled in the same stacking order, then rolled to a total thickness of 2mm non-composite laminate (NCL). Coating isolation agents are used to ensure that the interface does not bond after rolling so that the non-composite laminate (NCL) can be rolled in the same state as the laminate metal composite (LMC). Finally, the LMC and the NCL were simultaneously annealed at different processes (250°C/1h, 250°C/2h, 300°C/1h, 350°C/1h) under vacuum atmosphere to adjust the microstructure of the laminates and obtain heterogeneous lamella structure (HLS) Cu laminates. These four laminates are termed as HLS-a, HLS-b, HLS-c, and HLS-d. The name of samples and specific process procedures are listed in Table 2. The picture of LMC and NCL are shown in Fig. 2. It can be seen that the surface of the material was very smooth without defects, suggesting that the laminate processed good quality.

Table 2

Reduction ratio of constituent layers and post-annealing process in laminated Cu
Sample	Surface layer cumulative reduction	Center layer cumulative reduction	Post-annealing process
HLS-a	67%	87%	250℃/1h
HLS-b	67%	87%	250℃/2h
HLS-c	67%	87%	300℃/1h
HLS-d	67%	87%	350℃/1h

2.3 Characterization

Microstructural evaluation was investigated using optical microscope (OM), scan electron microscope (SEM) and electron backscatter diffraction (EBSD). Metallographic morphology was obtained by etching the laminated composite plate with aqueous ferric chloride hydrochloric acid solution (5 g FeCl₃ + 50 ml HCl + 100 ml H₂O), and the metallographic morphology of the TD plane was observed by optical microscopy. Fracture specimens obtained from tensile and tensile-shear tests were observed by scanning electron microscopy for fracture morphology. The EBSD specimens were prepared by electropolishing in a solution composed of 30% HNO₃ and 70% CH₃OH. The electropolishing voltage was 15 V. The scan step size was set as 0.5 µm to obtain EBSD mapping.

The test of Vickers microhardness was performed on the cross-section of samples using a load of 100 g for a dwell time of 10 s. To ensure reproducibility, the microhardness values of each sample were repeatedly tested at three separate positions. Dog-bone-shaped tensile specimens with a gauge length of 24 mm and a width of 4 mm were cut along the RD on pure copper laminates and non-compound triple copper plates. Tensile-shear specimens were cut along the RD on pure copper laminates. Tensile and tensile-shear specimens were tested on an INSTRON 5969 electronic universal testing machine at a stretching rate of 0.5 mm/min. During the tensile test, an extensometer was used and was removed when the strain reached 4%. As can be seen in Fig. 3, tensile tests consist of three parts: stretching the heterogeneous lamella structure (HLS) pure copper laminate specimens; stretching the surface layer and the center layer specimens of the non-compound laminate (NCL), respectively; stretching the non-compound laminate (NCL) specimens of combined stacking. To ensure data repetitiveness, at least 3 specimens were tested.

3.1 Microstructure characterization

Figure 4 are the optical microscopy images of all samples, which clearly show the interface areas. As a result of the difference between the accumulated reduction of the surface and center layers, the recrystallization onset temperatures on the surface and center layers are different, and the microstructures of these two layers are also different after annealing. After annealing at 250°C/1h, both the surface and center layers failed to completely recrystallize; the surface layer was fibrous, while the center layer consisted of fine equiaxed grains and fibrous tissue due to a greater degree of recrystallization than the surface layer. After annealing at 250°C/2h, with the extension of the holding time, equiaxed grains appeared in the surface layer, but still retained the fiber organization, and the number of equiaxed grains gradually increased in the center layer, but not uniformly. When annealing at 300°C/1h, with the increase in annealing temperature, the surface layer develops fine equiaxed grains, but still retains the original fiber organization. However, the center layer is recrystallized into equiaxed grains with the new grains in contact with one another, and the original fiber arrangement has been removed, indicating that recrystallization has been completed. When annealed at 350°C/1h, with the annealing temperature further increased, the surface layer consists of all new grains, the fibrous tissue disappears completely, and the center layer is composed of uniform equiaxed grains. It can be seen that the HLS-b and HLS-c exhibit a heterogeneous microstructure, with the surface layer composed of elongated grains and the center layer composed of approximately equiaxed grains.

Figure 4 (b) shows the microhardness of all samples along the ND direction. It can be observed that in the HLS-a, HLS-b and HLS-c, the hardnesses of the surface and center layers are clearly different, and there are obvious hardness incompatible interfaces at the bonding interface of adjacent layers. After annealing at 250°C/1h, the average hardness of the HLS-a was 94 HV for the surface layer and 83.3 HV for the center layer. Both the surface and center layers maintained high hardness levels due to the high work-hardening effect. After annealing at 250°C/2h, the average hardness of the surface layer was 83.56 HV and that of the center layer was 54.6 HV in the HLS-b, with significant differences between these two layers. After annealing at 300°C/1h, the average hardness of the surface layer was 73.85 HV and that of the center layer was 56.33 HV in the HLS-c. Compared with HLS-b, the hardness of the surface layer of HLS-c is slightly lower as a result of the higher recrystallization level of HLS-c than HLS-b, but the hardness of the surface and center layers of HLS-c are still significantly different. In the HLS-d, there was no significant change in hardness along the ND direction, with an average hardness of 54 HV. This is because both the surface and central layers have completed recrystallization so that they maintain the same hardness. In summary, the microhardness of the HLS-b and HLS-c exhibited a clear gradient distribution, which is very consistent with the unique microstructural characteristics of heterogeneous structural laminates. The apparent hardness difference indicates a significant mechanical incompatibility between the layers, which enhances their interaction during strain [13, 29].

3.2 Microstructure evolution

EBSD is an efficient way to analyze the microstructure characteristics. To unveil the microstructure evolution of all samples, the grain boundary characteristics and inverse pole figure (IPF) map of the surface and center layers are illustrated in Fig. 5 (a-d). One can see from Fig. 5 (a-d) that the interface between the constituent layers (indicated by white dotted line) was recognizable according to the difference in grain size. The colors of IPF map represent the crystallographic orientation of each grain parallel to the TD direction. Obviously, the grain orientation of the two layers under 250°C/1h, 250°C/2h and 300°C/1h annealing showed significant differences, and the grain orientation difference between the two layers decreased after 350°C/1h annealing treatment. The percentage of misorientation angle varies greatly with the change of annealing state in different samples. Figure 6 (a, b) shows that the LAGBs (2°–15°) and HAGBs (> 15°) distributions differs between the surface and center layers, indicating the heterogeneity of these laminated Cu. We found that the fraction of HAGBs in the surface layer under annealing at 250°C/1h is smaller than the fraction in the center layer, and that of HAGBs in the surface layer under annealing at 350°C/1h is comparable to that of the center layer. The fraction of HAGBs in the center layer increased significantly with 250°C/2h and 300°C/1h annealing compared to 250°C/1h annealing, but the change in the fraction of HAGBs in the surface layer was relatively small at the same annealing state. Therefore, after 250°C/2h and 300°C/1h annealing, a clear difference in misorientation angle can be seen between the surface and center layers. To evaluate the grain size, the average grain size (Fig. 6c) of the two layers of different samples was measured from the grain boundary feature map by Image J using the intercept method. The HLS-a exhibits elongated grains after annealing at 250℃ for 1h, but the grain width differs significantly. It is estimated that the average grain size in the ND direction of the surface layer is approximately 1.84 µm with an aspect ratio of 1:6.6, while the average grain size in the ND direction of the center layer is approximately 3.25 µm with an aspect ratio of 1:2.7, which indicates significant grain coarsening in the middle layer. After annealing at 250°C/2 h and 300°C/1h, the surface layers of both HLS-b and HLS-c exhibited elongated grains with an aspect ratio of about 1:3.3, while the center layers were both approximately equiaxed grains. After annealing at 350°C for 1 h, both the surface and center layers of the HLS-d had equiaxed grains, and the grain size of the center layer was slightly larger than that of the surface layer. There is a clear difference in grain structure between the surface layer and the center layer, which is attributed to higher recrystallization of the center layer.

To determine the degree of recrystallization, Fig. 7 (a-d) shows DefRex charts for all samples, the recrystallized grains are displayed in blue, deformed regions in red, and substructured (recrystallized with subgrains) grains in yellow. The DefRex plots demonstrate that the center layer recrystallizes more than the surface layer. The corresponding fractions of the surface and center layers in all samples shown in Fig. 7 (e,f). There is basically no recrystallization in the surface layer of HLS-a after annealing at 250°C/1h. Recrystallized grains are found in the center layer, but many deformed grains remain. After annealing at 250°C/2h or 300℃/1h, the degree of recrystallization of surface layer was low in HLS-b and HLS-c. In contrast, the recrystallized and substructured grains in the center layer occupied the whole area. After annealing at 350°C/1h, the surface and center layers of HLS-d crystallized to a high degree and the deformed grains disappeared. A greater degree of deformation of a metal typically results in more energy being stored and a stronger driving force for recrystallization. Center layers undergo a greater cumulative deformation than surface layers, so the center layers recrystallize more after annealing. Due to the distinct degree of recrystallization between the successive layers, there are obvious microstructural differences in grain orientation, misorientation angle and grain size between the surface and center layers. The resulting heterogeneous microstructure has a profound impact on the mechanical properties of pure copper laminate.

4.1 Bonding strength

Tensile-shear specimens were obtained from all samples, and their dimensions are shown schematically in Fig. 8 (a). The average interfacial shear strength of all samples was calculated by taking the average of three sets of test data, and the results are shown in Fig. 8 (b). The figure shows that the average interfacial shear strengths of the samples are 125.9 MPa for HLS-a, 131.02 MPa for HLS-b, and 178.5 MPa for HLS-c. The average interfacial shear strength of HLS-c was significantly higher than that of HLS-b. It shows that the 300°C/1h annealing treatment increases the elongation of the surface layer, reduces its deformation resistance, and enhances the interfacial shear strength of the laminate compared to the 250°C/2h annealing treatment. However, the interfacial shear strength of HLS-d reduced to 137.5 MPa with further increase in annealing temperature. This is due to recrystallization and softening of the laminate after annealing at 350°C/1h.

In order to further analyze the shear properties of the interface, the shear section morphology of all samples was measured. As shown in Fig. 8 (c-f), HLS-c and HLS-d have a greater amount of copper metal attached to the surface interface than HLS-a and HLS-b, and the surface area of the copper metal (shown in red) is also much larger. It indicates that with the increase in annealing temperature, the element diffusion ability is boosted and the laminate's interfacial bonding quality improves, as the shear strength of HLS-c prepared by annealing at 300°C/1h is also higher than that of HLS-b prepared by annealing at 250°C/2h. As the pure copper laminate is annealed at 350°C /1h, the copper laminate recrystallizes and softens, losing tensile strength and exhibiting mechanical properties that are similar to a single sheet of copper. Hence, HLS-d displays a low shear strength even with a well-bonded interface. By looking at the results, it can be seen that the HLS-c obtained by annealing at 300°C/1h achieves a strong interfacial bond and exhibits an ultra-high interfacial shear strength.

4.2 Tensile strength

The engineering stress-strain curves of all samples are shown in Fig. 9 (a), and the mechanical properties are shown in Fig. 9 (b). The good bonding of the interface supports the simultaneous fracture of the pure copper laminate. HLS-a is observed to have the highest strength and lowest plasticity after annealing at 250°C/1h, which is caused by the fact that the surface and center layers still retain the elongated grains formed after rolling and retain a lot of work hardening. After annealing at 350°C/1h, the strength of the pure copper laminate is significantly reduced and plasticity is significantly increased. HLS-d exhibits mechanical properties similar to a homogeneous copper plate because both the surface and center layers of the laminate are recrystallized and work hardening is completely eliminated after annealing treatment at 350°C/1h. After annealing at 250℃/2h or 300℃/1h, the strength of pure copper laminate decreases and plasticity increases significantly. HLS-b and HLS-c achieve a good strength-plasticity match because of the combination of a work-hardened surface layer and a completely work-hardening-free center layer. As compared to HLS-b, HLS-c showed similar elongation but significantly higher yield and tensile strengths. According to the results, the HLS-c annealed at 300°C/1h showed the best overall mechanical properties, with tensile strength and elongation at fracture reaching 278.08 MPa and 46.2%, respectively.

Figure 9 (c-f) displays the edge view of the fractured specimen and the enlarged views of the interface region for all samples. As you can see, the pure copper laminates prepared under 250°C annealing for 1h and 250°C annealing for 2h have relatively obvious interfacial separation and debonding, while the pure copper laminates prepared under 300°C annealing for 1h and 350°C annealing for 1h have no obvious interfacial separation, thus supporting a reliable interfacial bonding between different layers. HLS-a, HLS-b, and HLS-c fracture in shear mode in their thickness directions, and the shear surface is angled to the interface, indicating the stress-strain state generated by the interface interaction plays a large role in the process of tensile damage. As shown in Fig. 9 (g,h), in HLS-a and HLS-b, the surface and center layers were deformed independently after interface debonding, and the necking and dimples developed independently. Figure 9 (i) shows that the HLS-c prepared under annealing at 300°C/1h exhibits fracture characteristics completely different from those of the HLS-b prepared under annealing at 250°C/2h. A well-bonded interface exists between the layers of HLS-c during fracture, and the surface layer shows a smooth surface without dimples and only a shallow micropore. The center layer shows shallower microporosity and some oval depressions due to its higher ductility. Generally, microvoids in ductile material can grow and coalesce to large dimples only when the growth parameter (A) reaches the critical value (Ac) [30]:

$$A={\epsilon }_{p}\text{e}\text{x}\text{p}({\sigma }_{m}/{\sigma }_{0})$$

where ${\sigma }_{P}$ and ${\sigma }_{m}$ are accumulated effective plastic strain and triaxial stress level around microvoids, respectively; ${\sigma }_{0}$ is the yield strength. HLS-b and HLS-c have similar elongation, but HLS-c shows a higher yield strength. As a result of its lower growth parameters, HLS-c shows different fracture morphology characteristics than HLS-b. As shown in Fig. 9 (j), the different layers of HLS-d still maintain a good bonding after fracture, no obvious shear surface appears, and the intermediate layers present dense dimples, exhibiting the same fracture characteristics as a single copper plate. In summary, HLS-b and HLS-c exhibit different fracture characteristics due to different interfacial bonding conditions.

Both HLS-b prepared under annealing at 250°C/2h and HLS-c prepared under annealing at 300°C/1h exhibited good integrated mechanical properties, but HLS-c achieved the best integrated mechanical properties. To further analyze the difference between the heterogeneous laminates prepared under the two annealing processes, Fig. 10 (a) and Fig. 10 (b) show the stress-strain response of the heterogeneous lamella structure Cu laminates and the non-composite surface and center layers prepared under 250°C/2h annealing and 300°C/1h annealing, respectively. It should be emphasized that the surface and center layers of non-composite laminate undergo the same mechanical and heat treatment as the surface and center layers of pure copper laminate composite. The stress-strain response of HLS-b and HLS-c with their surface and center layers of non-composite plates is compared. According to the results, the surface and center layers exhibit completely different mechanical properties, which are caused by their distinctly different microstructures. The surface layer has low strain hardening because it retains a lot of work hardening and exhibits high strength with relatively low elongation [14], while the center layer has good elongation but exhibits low strength due to recrystallization softening. In comparison to the center layer, HLS-b and HLS-c exhibit a similar elongation but a much higher strength, resulting in an excellent combination of strength and ductility. This means that the combination of a hard layer (surface layer) and a soft layer (center layer) can effectively integrate strength and ductility. On the contrary, the three-layer non-composite layered specimen (NCL) under combined tensile manifests very low strength after fracture of the center layer, implying that the deeply bonded interface significantly contributes to the coordinated deformation during stretching.

According to the rule of mixing (ROM), yield strength, tensile strength, and uniform elongation were predicted for HLS-b and HLS-c. The ROM calculation formula is as follows [31]:

$${\sigma }_{ROM}={V}_{F}\bullet {\sigma }_{F}+{V}_{C}\bullet {\sigma }_{C}$$

Where ${\sigma }_{ROM}$,${\sigma }_{F}$ and ${\sigma }_{C}$ represent the flow strength of the composite laminates, the surface layer, the center layer, respectively, ${V}_{F}$ and ${V}_{C}$ are the cross-sectional area fraction of the surface layer (62.2%) and center layer (37.8%), respectively. The ultimate elongation of composite laminates is generally estimated by using the force based ROM [13]. The ultimate elongation can then be expressed as:

$${\epsilon }_{ROM.UE}=\frac{{V}_{F}{\sigma }_{F.UE}{\epsilon }_{F.UE}+{V}_{C}{\sigma }_{C.UE}{\epsilon }_{C.UE}}{{V}_{F}{\sigma }_{F.UE}+{V}_{C}{\sigma }_{C.UE}}$$

Where ${\epsilon }_{UE}$, ${\sigma }_{UE}$ represent the strain and stress of the corresponding component at the necking point. Figure 10 (c-e) shows the mechanical properties of laminates, surface and center layers annealed at 250℃/2h and 300℃/1h, and the values calculated using the rule of mixing. In the results, it is noted that the uniform elongation of both HLS-b and HLS-c is substantially higher than in the ROM calculation, which also demonstrates that the more ductile middle layer greatly contributes to the deformation of the less ductile outer layer. The yield and tensile strengths of HLS-b are lower than the ROM calculated values. However, the experimental values for the strength of HLS-c are similar to the ROM calculated values, which follows the prediction criterion of ROM. The yield, tensile, and ROM calculated values of the surface and center layers obtained after 300°C/1h annealing were lower than those obtained after 250°C/2h annealing, but the strength of HLS-c prepared at 300°C/1h annealing was significantly higher than that of HLS-b prepared at 250°C/2h annealing and yield strength increased by 26.6%, which can be attributed to the synergistic strengthening of the surface and center layers due to the good bonding of the interface during the tensile process.

In summary, the surface and center layers of the laminates prepared under 250°C/2h annealing and 300°C/1h annealing had obvious differences in mechanical properties due to their microstructural differences. The severe mechanical incompatibility of the surface and center layers caused synergistic constraints between the layers during the tensile process, which enabled the heterogeneous microstructured copper laminates to achieve good overall mechanical properties. The mechanical properties of both the surface and center layers of HLS-b are superior to those of HLS-c. HLS-c achieves the optimal combination of mechanical properties due to its high interfacial bonding strength (178.5 MPa). It indicates that a good bonding interface during the tensile process will enhance the synergistic confinement and strengthening between the different layers, allowing the material to exhibit excellent mechanical properties.

In this study, pure copper laminates with heterogeneous microstructure were fabricated by rolling bonding with subsequent heat treatment, and the microstructure and mechanical properties of the laminates were characterized. The main conclusions are as follows:

In this paper, a high-efficiency and scalable method of preparing heterogeneous lamella structure (HLS) Cu laminates was proposed. HLS with different cumulative reduction of the surface and center layers were prepared by cold-rolling pretreatment of the center layer (60% reduction) and cold-rolling bonding of the whole blank (67% reduction), and the heterogeneous microstructure was optimized by a subsequent annealing process. For comparison, the non-composite laminate (NCL) was fabricated by performing the same rolling and annealing treatments as the HLS, except that a separator was added between the surface layer and the center layer in the blank. After rolling and annealing, the surface layer and center layer could be separated from the NCL.

The results of the EBSD analysis show that the degrees of recrystallization of the surface and center layers of the HLS are different, due to the different accumulated deformation in the surface and center layers of the HLS. Therefore, the surface and center layers show different microstructures and possess different mechanical properties, which promotes the achievement of significant strengthening and strain hardening.

The results of the tensile test on HLS and those of the surface layer and the center layer separated from the NCL indicate that HLS at 300°C/1h exhibits excellent comprehensive mechanical properties. It showed excellent elongation, which was higher than the calculated values based on ROM, and exhibited significantly greater strength than the fully annealed center layer. The yield strength, tensile strength and elongation at break reaching of the HLS specimens are 139.13 MPa, 278.08 MPa and 46.2%, respectively. The HLS annealed at 300°C/1h exhibited the highest interfacial bonding strength of 178.5 MPa. The tensile results of NCL specimens and the results of tensile shear tests show that the deeply bonded interface significantly contributes to the coordinated deformation during stretching, it has a positive effect on the improvement of the mechanical properties of HLS.

The difference in mechanical properties of the heterogeneous layers in the HLS and the bonding state of the interface together influence the improvement of the overall performance of the material. During the tensile process, the inhomogeneous deformation of the heterogeneous layers leads to synergistic strengthening and strain hardening of the material, and a good bonding interface optimizes the synergistic constraint between the heterogeneous layers, which effectively enhances the comprehensive mechanical properties of the laminate and promotes it to achieve a good strength-plasticity match.

Acknowledgements

Not applicable

Funding

Supported by National Key Research and Development Program (2018YFA0707300), General Program of National Natural Science Foundation of China (grant number 51905372), and postdoctoral Science Foundation of China (grant number 2020T130463).

Availability of data and materials

The datasets supporting the conclusions of this article are included within the article.

Authors’ contributions

The author’ contributions are as follows: Zhi-Hui Gao wrote the manuscript; Tao Wang proposed ideas and edited revised the manuscript; Yun-Lai Zhao revised the manuscript; Hua Ding assisted with the language of the manuscript; Qing-Xue Huang gave some advice on the manuscript. All authors read and approved the final manuscript.

Competing interests

The authors declare no competing financial interests.

Consent for publication

Not applicable

Ethics approval and consent to participate

Not applicable

Biographical notes

Zhi-Hui Gao, born in 1997, is currently a master candidate at Engineering Research Center of Advanced Metal Composites Forming Technology and Equipment of Ministry of Education, Taiyuan University of Technology, China.

Tel: +86-18406583713; E-mail: [email protected]

Tao Wang, born in 1985, is currently a professor and a PhD candidate supervisor at Engineering Research Center of Advanced Metal Composites Forming Technology and Equipment of Ministry of Education, Taiyuan University of Technology, China.

Tel: +86-18735170186; E-mail: [email protected]

Yun-Lai Zhao, is currently a Ph.D. candidate at School of Mechanical and Aerospace Engineering, Jilin University, China

Tel: +86-18004427876; E-mail: [email protected]

Hua Ding, born in 1979, is currently a professor and a PhD candidate supervisor at Shanxi Key Laboratory of Fully Mechanized Coal Mining Equipment, Taiyuan University of Technology, China.

Tel: +86-18835123666; E-mail: [email protected]

Qing-Xue Huang, born in 1960, is currently a professor and a PhD candidate supervisor at Engineering Research Center of Advanced Metal Composites Forming Technology and Equipment of Ministry of Education, Taiyuan University of Technology, China.

E-mail: [email protected]

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Editorial decision: Minor revision
10 Jan, 2023
Reviewers agreed at journal
07 Jul, 2022
Reviewers invited by journal
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Editor invited by journal
22 Jun, 2022
Editor assigned by journal
06 Jun, 2022
First submitted to journal
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Mechanical properties optimization and microstructure analysis of pure copper gradient laminates via rolling

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Abstract

Figures

1. Introduction

2. Materials And Methods

2.1 Materials Preparation

2.2 Preparation processing

2.3 Characterization

3 Heterogeneous Structure

3.1 Microstructure characterization

3.2 Microstructure evolution

4 Performance Testing

4.1 Bonding strength

4.2 Tensile strength

5 Discussion

6 Conclusions

Declarations

References

Supplementary Files

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