Enhanced strength-ductility synergy in an ultra-strong copper alloy via coherent nanoprecipitates and stress-induced twinning

Deformation twins (DTs) and stacking faults (SFs) are rarely found in Cu-20Ni-20Mn alloy, which limits further strengthening of alloy. Here, we report ultra-strong Cu-20Ni-20Mn alloy prepared by introducing a high density of coherent nanoprecipitates. The ultimate tensile strength (UTS) of alloy can be largely enhanced to 1624.7 ± 14.7 MPa, while maintaining an acceptable elongation of 7.3 ± 0.5%. We find that the coherent nanoprecipitates contribute to an ultra-high true stress up to about 1740 MPa, which facilitate the formation of DTs. The strength and ductility improvement stems from the synergistic effect of dislocations, nanoprecipitates and DTs. GRAPHICAL ABSTRACT IMPACT STATEMENT The Cu-20Ni-20Mn alloy sheet exhibiting an unprecedented strength-ductility synergy has been prepared by a two-step rolling process. The ultrafine grain structure containing substantial deformation twins coupling with homogeneously distributed nano-coherent precipitates accounts for the superior mechanical properties.


Introduction
Structural materials with simultaneously ultrahigh strength and good ductility are often required for industrial applications.However, strength and ductility are a couple of contradictions in most materials [1][2][3][4][5][6][7].It is well known that traditional strengthening methods such as solid solution strengthening, grain refinement and precipitation strengthening often lead to reduced ductility or brittleness due to dislocations accumulating at internal defects and boundaries [4,[8][9][10][11][12].The internal boundaries introduced by the above-mentioned traditional strengthening methods are incoherent.The incoherent boundaries can effectively restrict dislocation movement to enhance strength, but the capacity to adapt to plastic deformation is impaired by reducing ductility and toughness [1,13].Previous studies have shown that introducing stable and nano-coherent internal boundaries offers the possibility for considerable strengthening while retaining an appropriate range of ductility [14][15][16][17].Coherent interfaces can be created by the introduction of twins, partial dislocations and coherent nanoprecipitates [18].
Copper alloys with high strength and good electrical conductivity are widely used as components in aerospace industry, medical devices, marine engineering, automotive parts and electronic appliances [19][20][21][22].Among the wide variety of copper alloys, the Cu-20Ni-20Mn alloy shows excellent mechanical properties, which has a tensile strength of 1200-1400 MPa and an elastic modulus around 150 GPa.However, the elongation of this ultra-strong alloy is only 2-4% which brings risks in application if the brittle fracture happens [23][24][25][26].The main strengthening mechanism for the Cu-20Ni-20Mn alloy is precipitation strengthening due to its high stacking fault energy (SFE) [26].Precipitation-hardening as an available strain-hardening mechanisms of this alloy remains confined to the interactions of dislocations, precipitates and grain boundaries, which often reduces ductility [27][28][29][30].As a result, this alloy has so far not released the considerable strain-hardening capacity provided by stacking faults (SFs) and twins.In addition, a lamellar discontinuous NiMn phase forms during aging treatment [26,31].Pile-up of dislocations surrounding the discontinuous precipitates can lead to severe strain localization at the interfaces which are mismatched with the matrix lattice, finally resulting in cracks and catastrophic failure [12,[32][33][34][35]. Like in many other structural materials, it remains a long-standing challenge how to strengthen the Cu-20Ni-20Mn alloy at a minimized sacrifice of ductility [13].Introducing coherent, stable internal boundaries may offer the possibility for synergistically improve strength and ductility, as demonstrated in nanocrystalline and nanotwinned materials [17,36,37].Coherent interfaces develop in the twin boundaries or interfaces of coherent nanoprecipitates.However, deformation twins (DTs) and SFs have not been reported yet for tensile loaded Cu-20Ni-20Mn, as the deformation mechanism is dominated by dislocation slip which is less energy costly [38,39].Recently, Wang et al. [40] revealed that nanoprecipitates contribute to an ultrahigh strengthening effect which can enhance the true tensile stress to reach critical value for the formation of mechanical twins in a compositionally complex steel, in turn, enhances the mechanical property by further improving strain hardening capacity and toughening.
Inspired by these studies and in order to design copper alloys with high strength and ductility, we report here on DTs and the corresponding strengthening effect in the Cu-20Ni-20Mn alloys with high SFE.That is to say, the coherent nanoprecipitates are obstacles and sources of dislocations through dislocation cutting and bypassing, producing high strain hardening.In addition, at ultrahigh stress and strain level, which results from the interactions between dislocations and coherent a large amount of nanoprecipitates and the increasing dislocation densities, activate the previously unattainable twinning mechanism by deformation-induced in the Cu-20Ni-20Mn alloy.The twinning-induced plasticity effect is an efficient toughening and strain hardening mechanism, enables an excellent strength and ductility synergy [41][42][43].In this work, we report a way to introduce coherent interfaces in the Cu-20Ni-20Mn alloy, through trace addition of Nb and employing a two-step rolling process.Thereby, a large number of coherent nanoprecipitates and DTs can be created in Cu-20Ni-20Mn-xNb (x = 0.4, 0.8 wt.%) alloys, providing a good strengthductility balance.The proposed approach provides a new strategy to achieve ultrahigh strength with good ductility for advanced copper alloys.

Materials and methods
In this work, three Cu-20Ni-20Mn-xNb alloys (x = 0, 0.4, 0.8 wt.%) (referred to as 0Nb, 0.4Nb and 0.8Nb) were chosen as target materials.Dog-bone-shaped tensile samples were cut along the rolling directions.Tensile tests were carried out at room temperature with a strain rate of 1 × 10 −3 s −1 using a SHIMADZU-AG-Xplus 50 kN electronic universal testing machine equipped with an extensometer.Nanoindentation test was conducted using a KLA-iMicro Nano Indenter with a peak force of 50 mN and a holding time of 10 s.Indentation matrices (15 × 15) with a spacing of 10 μm between adjacent indents were conducted on each sample.The specific preparation method and microstructure characterization of materials is shown in the Supplementary material.

Results and discussion
Representative tensile engineeringstress-strain curves of the 0Nb, 0.4Nb and 0.8Nb alloys after solution treatment (referred to as the ST-0Nb, ST-0.4Nb and ST-0.8Nb samples hereafter) and aging treatment (referred to as AT-0Nb, AT-0.4Nb and AT-0.8Nb samples hereafter) are depicted in Figure 1(a) (the corresponding tensile fracture morphology is shown in Figure S1).After solution treatment, the Nb-including sample has higher ductility but lower strength than the Nb-free sample.After aging treatment, the strength of alloys increased but the ductility decreased on account of precipitation hardening, compared with the AT-0Nb sample, the strength and ductility of the AT-0.4Nb and AT-0.8Nb samples are significantly improved.The ultimate tensile strength (UTS) and yield strength (YS) of the AT-0.4Nbsamples reach ∼ 1624 MPa and ∼ 1463 MPa, respectively, which are increased by ∼ 126% and ∼ 236% from those of the ST-0.4Nbsamples.Notably, the YS and UTS of the AT-0.8Nbsamples do not continue to increase, but decrease,  [19,47] and Cu-Be [20,48], Cu-Ti [45,49], Cu-Ni-Si [50] and Cu-Ni-Mn [24,25,51] alloys.
due to the formation of greater number of Ni 3 Nb particles in the matrix, the volume fraction of Ni 3 Nb particles in 0.4Nb sample is about 1.5%, and the volume fraction of Ni 3 Nb particles in 0.8Nb sample is about 3%.Moreover, the Ni 3 Nb particles are easily enriched and continuously distributed on the recrystallized grain boundaries (Figure S6(i-k)), and more Ni 3 Nb particles increase the possibility of grain boundary fracture, which reduces the strength and ductility of the alloy [23,44].The corresponding true stress-strain curve are depicted in Figure 1(b), revealing a super high true stress up to ∼ 1740MPa for the AT-0.4Nbsamples.The aged Nbincluding samples exhibit improved strength and ductility of the Nb-free sample.The strain hardening rate (Figure 1(c)) of AT-0.4Nb is evidently higher than for the other two types of samples at any given strain, contributing to the delay of plastic instability.Moreover, when comparing the UTS and elongation of our ATsamples with previously reported high strength and high elastic copper alloys in Figure 1(d) shows that our AT-0.4Nballoy exhibits an excellent combination of UTS and uniform elongation, which has significant engineering application value [19,20,25,45,46].
Figure 2 shows the microstructure of the different samples after aging treatment.The electron backscatter diffraction (EBSD) results of Cu-20Ni-20Mn-xNb alloys after aging treatment are shown in Figure 2(a-f).A large number of deformed elongated grains are visible for the AT-0Nb samples (Figure 2(a)), and the fraction of deformed elongated grains decreases with increasing Nb content (Figure 2(a-c)).The recrystallized microstructure is comprised of uniformly distributed equiaxed grains with an average size of ∼ 2.5 μm (Figure S2(a-c)).The distribution of recrystallized grains in the samples with different Nb content is shown in Figure 2(d-f).For the AT-0Nb samples, the thermal activation energy provided by the solution temperature is not sufficient for complete dynamic recrystallization.Only some dynamic recrystallization occurs at grain boundaris/dislocation networks, and a large number of severely deformed grains (non-crystallized) still exists, so that the volume fraction of dynamically recrystallized grains is only ∼ 36.74%.The volume fractions of dynamically recrystallized grains in the AT-0.4Nb and AT-0.8Nb samples are ∼ 58.75% and ∼ 80.96%, respectively, thus increasing with increasing Nb addition.Both sub-grain (substructured, i.e. recovered grain) boundaries and highangle grain boundaries (HAGBs) (Figure S2(d-f)) can absorb dislocations, reducing the dislocation density of the matrix and improving the ability to store dislocations [52] The Kernel Average Misorientation (KAM) and GOS values (Figure S2(g-i) and Figure S3) are reduced with the addition of Nb, that is, the dislocation density and lattice distortion are reduced, which improves the uniform deformation ability of the alloy.The inverse pole figures of the samples with different Nb contents are depicted in Figure 2(g).The AT-0Nb alloy exhibits a strong < 101 > texture along the rolling direction.From the results, it can be concluded that Nb addition not only promotes complete recrystallization but also weakens the < 101 > texture, which will facilitate homogenous deformation, resulting in a more superior ductility of the AT-0.4Nb and AT-0.8Nb samples.These findings reveal that Nb promotes recrystallization, and improves the strength and ductility of the Cu-20Ni-20Mn alloy.As shown in Figure 2(h), the aged Cu-20Ni-20Mn-xNb alloys exhibit a face centered cubic (FCC) structure.In addition to the main FCC α-Cu diffraction peaks, some secondary peaks are found for the AT-0.4Nb and AT-0.8Nb samples, whereas they are absent in the case of the ST-0Nb and AT-0Nb samples.The secondary peaks are determined to correspond to the Ni 3 Nb phase.The diffraction peaks of the FCC α-Cu phase are broader and shifted significantly to higher diffraction angles (2θ) after aging treatment, which indicated that a large number of face-centered tetragonal (FCT) NiMn precipitation are precipitated from the α-Cu matrix during aging.
Figure 3 shows the microstructure evolution of the fractured Cu-20Ni-20Mn-xNb alloys after tension.Figure 3(a-d) present the microstructure for the AT-0Nb samples.The yellow line of Figure 3(a) reveals that the NiMn phase of the AT-0Nb samples is lamellar, and its corresponding energy dispersive spectroscopy (EDS) line scans across Cu / discontinuous precipitations is shown in Figure S4.A large number of lamellar NiMn phases and the Ni-rich phases which hinder the movement of dislocations (Figure S5).Dislocation pile-up leads to stress concentration and deterioration of the mechanical properties of the alloy.Figure 3(c) and Figure S6(a-d) show that the slip mode of the dislocation is wavy slip, and very high densities of dislocations cause stress concentration of the alloy producing micro-cracks, resulting in failure of the material.The microstructure of the AT-0.4Nbsamples are presented in Figure 3(e).The addition of Nb inhibits discontinuous precipitation (DP) of particles and transforms them into uniformly dispersed nanoprecipitates.Figure 3(f) displays a high-resolution transmission electron microscope (HRTEM) image of the nanoprecipitates in Figure 3(e).The fast Fourier transform (FFT) pattern corresponding to Figure 3(f) shows two set of diffraction spots (Figure 3(g1)).No misfit dislocations were observed at the interface between the nanoprecipitates and the α-Cu matrix, as shown by inverse FFT (IFFT) images (Figure 3(g2-g4)).It is obvious that the nanoprecipitates are coherent with the α-Cu matrix [32].Many nanoprecipitates interact with the dislocations lead to high strengthening effect during deformation, which increases the flow stress to approach critical value for the formation of DTs and SFs (Figure S7), generating a large number of DTs (Figure S8), thereby improving the storage capacity of dislocations, endowing alloys with outstanding strain hardening capacity [39,53].The strength and ductility of the alloy are improved simultaneously through the interaction of precipitates and DTs with the dislocations.With further increasing Nb content, a number of Ni 3 Nb particles and nanoprecipitates are formed (Figure 3(i,j)).The Ni 3 Nb particles hinder the movement of dislocations and dislocation pile-up leads to stress concentration (Figure S6(i-k)), which reduces the strength and ductility of the alloy compared with the 0.4Nb samples, which is confirmed by the corresponding fracture morphology.A large number of DTs are also observed in the material (Figure 3(k,l) and Figure S9(a-c)), which improves the strain hardening ability and uniform deformation ability of the alloy.
The microstructure inhomogeneity of adjacent domains is caused by DP and the Ni-rich phase can readily lead to local deformation due to the strength mismatch of the different domains.It is speculated that the addition of Nb strongly alleviates the inhomogeneity of the microstructure and promotes the uniform deformation of the alloy.This hypothesis is verified by nanoindentation tests, and the test results are shown in Figure 4.With the increase of Nb content, the hardness and modulus of the alloys decrease slightly, as shown in Figure 4(c,d).
In different regions of the AT-0Nb samples, both hardness and modulus show a remarkable intensity variation, which may be due to the obvious difference in the elastic-plastic deformation resistance of the different areas caused by DP and the Ni-rich phase.In contrast, the changes in hardness or modulus in the AT-0.4Nbs and AT-0.8Nb samples are much smaller in the different domains, as shown in Figure 4(c).This indicates that stress matching of the adjacent domains after the decoration with dispersed coherent nanoprecipitates reaches a close level through precipitation hardening.Namely, the dispersed coherent nanoprecipitates exert rigid deformation constraints on the ductile α-Cu matrix and promote a coordinated deformation of the different domains [54][55][56].Compared with the AT-0Nb samples, the distribution of hardness and modulus in the AT-0.4Nbsamples is more homogeneous, which confirms that dispersed coherent nanoprecipitates are conducive to strength matching of different domains, thus providing compatible deformation capability in different domains [57][58][59] and contributing to the uniform plastic deformation of the alloy.
Based on the TEM results and deformation microstructures, a schematic diagram describing the deformation mechanism of the Cu-20Ni-20Mn-xNb alloys can be proposed, as shown in Figure 5. Upon deformation, dislocations accumulate at incoherent DP and Ni-rich phase precipitates, which causes an uneven stress distribution and microcracking, thus driving the AT-0Nb samples to fracture.The AT-0.4Nb and AT-0.8Nb samples at high stress levels can activate previously unattainable deformation induced twinning mechanisms in the FCC Cu-20Ni-20Mn alloy due to complex interactions between dislocations and nanoprecipitates.The formation of DTs, further improving toughening reserves  and strain hardening capacity duo to dislocation storage and pinning by DTs in the later stages of deformation, thus counteracting strain localization and softening, and improving the strength and ductility of the alloy [60][61][62][63].The dislocations planar slip and twin boundaries can effectively improve the strain hardening and uniform plastic deformation ability, and then causing an excellent strength-ductility match in the Cu-20Ni-20Mn alloys.Accordingly, the results show that the introduction of DTs through precipitation strengthening enhances the flow stress to approach the high critical value for the onset of mechanical twinning in high SFE Cu-20Ni-20Mn alloys, which can effectively bring an excellent strength-ductility synergy.

Conclusion
In the current Cu-20Ni-20Mn-xNb alloys the significantly enhanced strain hardening and tensile strength of the α-Cu matrix, resulting in the formation of DTs, together with the co-deformation via dislocations and coherent nanoprecipitates, helps to avoid stress localization.The addition of a proper amount of Nb strongly alleviates the mismatch of the copper matrix and promotes a uniform deformation of the alloy.Furthermore, the formation of DTs can effectively inhibit the nucleation and propagation of microcracks, resulting in remarkable toughening effect.Therefore, good tensile elongation ( ∼ 7.3%) is achieved, even at such an ultrahigh tensile strength ( ∼ 1624 MPa) in our Cu-20Ni-20Mn-0.4Nballoy.In summary, our design strategy provides new insights to enhance the mechanical properties of high SFE high-performance materials by triggering a previously unattainable deformation mechanism.

Disclosure statement
No potential conflict of interest was reported by the author(s).

Figure 4 .
Figure 4. Nanoindentation results of the Cu-20Ni-20Mn-xNb alloys after aging treatment.(a) Indentation map of the performed nanoindentation tests.(b) Load-depth curves.(c) Hardness and modulus distribution maps.(d) Frequency histograms of hardness and modulus.

Figure 5 .
Figure 5. Schematic illustration of the deformation mechanisms of the Cu-20Ni-20Mn-xNb alloys during tension testing.