Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... 415–422 MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF NANOMETER MULTILAYER Cu/Ta MATERIALS MOLEKULARNO DINAMSKE SIMULACIJE NATEZNO DEFORMIRANIH NANOMETERSKIH VE^PLASTNIH MATERIALOV NA OSNOVI BAKRA IN TANTALA Hu Jianian 1, 2 , Zhang Haotian 2 , Zhu Youlin 3 , Li Peibo 3 , Luo Guoqiang 1,3,* , Wang Chuanbin 1,3 , Shen Qiang 1,3 , Zhang Lianmeng 1,3 1 Chaozhou Branch of Chemistry and Chemical Engineering Guangdong Laboratory, Guangdong, Chaozhou, China 2 State Key Laboratory of Precision Blasting, Jianghan University, Hubei, Wuhan, China 3 State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Hubei, Wuhan, China Prejem rokopisa – received: 2022-05-26; sprejem za objavo – accepted for publication: 2022-07-04 doi:10.17222/mit.2022.505 In this research, the tensile mechanical properties and microstructure evolution of Cu/Ta nanolayered composites were studied using the molecular dynamics simulation method. By analyzing the tensile stress/strain relationship of Cu/Ta with different in- terface structures and the movement of dislocations during the stretching process, the deformation mechanism of materials with different interface structures and the effect of interface structures on the tensile strength of Cu/Ta nanolayered composites are revealed. The effect of shear localization during extension is also analyzed. The results show that the dislocation structures at the interfaces of Kurdjumov-Sachs-type and Nishiyama-Wasserman-type samples are parallelogram and triangular interface de- fect arrays, respectively, which can easily induce two Shockley partial dislocations to slide along different (111) planes, forming an intersection and merging into ladder-rod dislocations. However, dislocations between the Kurdjumov-Sachs {112}-type sam- ple interfaces exhibit parallel array characteristics, while the interfacial dislocations have non-planar interface components, which can induce deformation twinning. The process is dissociated through a set of intrinsic interfacial dislocations. Shockley partial dislocations are then formed by dislocation motion, creating stacking faults (SF1), and then the second set of partial dislo- cations may nucleate from the interface and slide on the adjacent SF1 plane, eventually forming deformation twinning. Keywords: molecular dynamics simulation, Cu/Ta nanolayered composites, tensile properties, deformation mechanism, disloca- tion motion Avtorji opisujejo natezne mehanske lastnosti in razvoj mikrostrukture ve~plastnih Cu/Ta nano-kompozitov s pomo~jo metod molekularno-dinamske simulacije. Z analiziranjem zveze med natezno napetostjo in deformacijo Cu/Ta z razli~no strukturo mejnih ploskev in gibanjem dislokacij med procesom raztezanja so avtorji odkrili kak{en je deformacijski mehanizem materialov z razli~no strukturo mejnih ploskev in njihov vpliv na natezno trdnost ve~plastnih Cu/Ta nano-kompozitov. Ugotavljali so kak{en je deformacijski mehanizem in vpliv lokaliziranega striga med raztezanjem. Rezultati raziskave ka`ejo na to, da so strukture vzorcev dislokacij na mejnih ploskvah Kurdjumov-Sachsovega in Nishiyama-Wassermanovega tipa s tri- in {tirikotno ureditvijo napak na mejnih ploskvah. To lahko enostavno spro`i (inducira) dve Shockleyjevi parcialni dislokaciji, ki zdrsita vzdol` razli~nih (111) ravnin in tvorita prese~i{~e dislokacij z obliko zlo`ljivih lestev. Vendar vzorci dislokacij Kurdjumov-Sachsovega {112}-tipa na mejnih ploskvah ka`ejo zna~ilno vzporedno razporeditev. Pri tem imajo mejne dislokacije neplanarne (neravninske oz. 3d) komponente, ki lahko inducirajo deformacijo z dvoj~enjem. Proces poteka z razdru`evanjem vrste notranjih mejnih dislokacij. Pri tem nastajajo Shockleyjeve parcialne dislokacije, ki tvorijo napake zloga (SF1). Nato lahko nastane (nukleira) {e drugi set parcialnih dislokacij na mejnih ploskvah in zdrs na sosednje SF1 ravnine, kar lahko privede do deformacije z dvoj~enjem. Klju~ne besede: molekularno-dinamske simulacije, ve~plastni nanokompoziti Cu/Ta, natezna deformacija, mehanske lastnosti, gibanje dislokacij 1 INTRODUCTION Tantalum has excellent properties, such as high melt- ing point, high strength and good wear resistance, while copper exhibits high thermal conductivity, excellent me- chanical properties and good electrical conductivity. 1–4 Therefore, in recent years, nanometer Cu/Ta multilayer related reports on composite materials (MNCs) have be- come a hotspot of academic research. 5,6 Copper/tantalum nanocomposites combine the advantages of Ta and Cu and have a wide range of applications in the electronics industry, used for semiconductor and superconductor fabrication, optical and magnetic systems, etc. Combina- tions of PVD, cross accumulation rolling (CARB) and intermediate annealing steps are the most commonly used techniques for the preparation of nanostructured Cu/Ta MNCs and bulk Cu/Ta nanolayered multilayers. 7–9 However, the nanostructure of Cu/Ta MNCs prepared with these methods is not uniform, the thickness is un- controllable and the preparation cost is high. 10 With the Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 415 UDK 539.382:669.294:669.3 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 56(4)415(2022) *Corresponding author's e-mail: hujianian1@outlook.com (Hu Jianian) rapid development of simulation technology, the use of molecular dynamics simulation technology to calculate and optimize the parameters of the nano/macrostructure and discover the evolution mechanism of the micro- structure during the deformation process has become an important tool for the study of metal nanocomposites. 11 Grain boundaries of MNCs are important for the ma- terial properties such as strength, fracture, work harden- ing and damage evolution under irradiation and im- pact. 12,13 R. Bejaud et al. 14 revealed that the Cu-Ag inter- face directly or indirectly induces the nucleation of twinning dislocations through Lomer dislocations, and through an atomic-scale analysis, they revealed that twinning is a common plasticity mechanism. Lu et al. 15 studied the effect of the Al/Cu MNC interface stress on the dislocation nucleation and propagation law during deformation and established a model of the interface stress, dislocation nucleation and motion parameters in an analytical form. The study shows that the yield and failure of Al/Cu MNCs are mainly controlled by the evo- lution of dislocations in the Al and Cu layers, respec- tively. Due to the interactions between stacking faults and dislocations, Al(100)/Cu(100), the shear strength of the MNCs decreases with decreasing the repeating inter- layer spacing. Interfacial microstructure characteristics of MNCs are important for their mechanical properties, such as in- terfacial defects, modulation period, interfacial structure, dislocation nucleation and deformation twinning nucle- ation. 16 Lu et al. 17 used molecular dynamics (MD) simu- lations to investigate the effect of interfacial structures such as interfacial defects and modulation period ( )on the deformation mechanism of Cu/Ta nanometallic multilayers (MNCs) under uniaxial stretching. The re- sults show that the Cu(111)/Ta(110) interface is the ini- tial site of dislocation nucleation, hindering further diffu- sion of dislocation motion toward the interface. During the tensile process involving Cu/Ta MNCs, the maxi- mum stress decreases with the increase in , and there is a critical point where the flow strength of the MNCs is the largest. However, the effects of interfacial dislocation nucleation and deformation twinning nucleation on the mechanical properties of Cu/Ta MNCs are lacking. Interfacial dislocation nucleation and deformation twinning nucleation are closely related to the interfacial orientation of nanolayered multilayer films. 18 Zhang et al. 19 investigated the nucleation of lattice dislocations at the Cu/Nb interface with the Kurdjumov-Sachs (KS) ori- entation relationship using atomic-scale simulations, sys- tematically explored the possibilities for the slip system, and concluded that, contrary to the initial design, the lo- cation nucleation sites of misfits correlate poorly with pre-existing misfit dislocations. Beyerlein et al. 20 investi- gated the effect of the intrinsic bimetallic interface struc- ture on lattice dislocation nucleation. It was found that the interface structure has a great influence on disloca- tion nucleation and the type of nucleation sites. At the same time, it was found that interfacial dislocations with non-planar Burgers vectors can prevent interfacial slid- ing, and the interface is more prone to twinning nucle- ation. Therefore, in this research, the MD method was used to study the influence of the interface structure of Cu/Ta materials with different orientation relationships on the mechanical properties. Interface structure characteristics of the three interface orientations were compared, and the effects of the interface structure and stretching direc- tion on the deformation mechanism were discussed. In this research, we studied the deformation mechanism of Cu/Ta nanolayered composites, emphasizing how the in- terface structure affects the interaction with defects, the strengthening mechanisms of samples with different in- terface structures and the effect of the loading direction on the mechanical properties of the samples. 2 COMPUTATIONAL DETAILS 2.1 General We selected three kinds of Cu/Ta nanolayered com- posites with classical interface structure orientations. The interface structures included the Kurdjumov-Sachs type (KS), Nishiyama-Wasserman type (NW) and Kurdjumov-Sachs {112} type (KS{112}). Their crystal orientations and box sizes are shown in Table 1. The molecular dynamics simulation used in this research was LAMMPS. The simulation process first used the conju- gate gradient (CG) method to obtain the minimum equi- librium energy model. The model optimization-process energy and force-pause tolerances were 10 –7 eV/nm and 10 –13 eV/nm, respectively. The simulation calculations of the tensile deformation of Cu/Ta MNCs were performed within the NpT ensemble, the temperature was con- trolled by the Nose-Hoover system, and the pressure was controlled by uniaxial stretching of the box along the z-axis direction. The calculation results of MD were post-processed by the OVITO visualization software. By identifying the crystal lattice types and structural defects in atomic crystals, the microstructure evolution process Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... 416 Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 Table 1: Cu/Ta MNCs with Kurdjumov-Sachs, Nishiyama-Wasserman and KS {112} interface orientations Interface KS NW KS{112} Constituent Cu layer Ta layer Cu layer Ta layer Cu layer Ta layer Lx 114 [01–1 ] 102 [–11–1 ] 115 [01–1 ] 89 [001 ] 114 [–110 ] 102 [11–1 ] Ly 64 [–1211 ] 35 [–2 –1 1] 58 [–211 ] 55 [1–10 ] 44 [11–1 ] 35 [1–10 ] Lz 12 [111 ] 11 [011 ] 12 [111 ] 11 [110][ 112][ 112 ] and deformation mechanism of the Cu/Ta MNCs under tensile deformation were analyzed. 21 2.2 Interfacial structure characteristics Figure 1 shows side views of the interface structures after the relaxation equilibrium of the three models. Atomic lattice types were identified using the common neighbor analysis algorithm (CNA) in OVITO. In this figure, green, blue and red areas stand for FCC, BCC and HCP structural coordination atoms, respectively, while white ones stand for the "unknown or non-standard" structural atoms caused by high lattice distortion. It can be seen that all three model interfaces are sharp. The in- terfaces of the KS and NW samples are flat. In contrast, the interface of the KS{112}-type sample is non-planar, containing regular facets (white atoms in Figure 1c). At the same time, parallel out-of-plane stacking faults (red atoms in Figure 1c) are observed in the KS{112}-type sample. The interface organization structures of the three in- terface models are shown in Figure 2. They were identi- fied using polyhedral template matching (PTM) in OVITO. The results show that the interface defects of the KS sample form periodic parallelogram arrays, and that the defect regions (red HCP atoms) are located at the four corners of the parallelogram. The NW-type sam- ple-interface defect array is parallel and rod-shaped along the [01–1 ] direction, and the whole is a periodic triangle. The KS{112}-type sample interface contains parallel and discontinuous HCP and FCC coordination lines. The dislocations at the interface cannot be identi- fied by the dislocation extraction algorithm (DXA) so that the Thompson tetrahedron is required to identify them. The results show that the KS-type sample interface contains two sets of b1 = 1/2[01–1 ] and b2 = 1/2[10–1 ] types of total dislocations along the CA and CB trajecto- ries, respectively (Figure 2a). The interfacial dislocation type of the NW-type sample is a partial dislocation, and the Burgers vector is [2–1–1 ] (shown as b3 in Figure 2b). The interface of the in-plane component of the misfit dis- location of the KS{112}-type sample is b4 = 1/2[1–10 ] (shown as b4 in Figure 2c), and the out-of-plane compo- nent isb=1/6 [112 ] (Figure 1b). The interfacial struc- tures of the three models have different characteristics that affect dislocation nucleation and deformation during their stretching. 3 RESULTS Figure 3 shows the stress-strain curves of the Cu/Ta nanolayered composites with different interface struc- tures under the tensile load at 300 K. It can be seen from the figure that the stress increases linearly with the strain Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 417 Figure 2: Atomic structures of the interfaces identified with PTM: a) KS, b) NW, c) KS{112} Figure 1: Geometric models of the samples after relaxation: a) KS type, b) NW type, c) KS {112} type until the first peak, which is the yield point of the Cu layer. Then the stress decreases over a small segment and begins to increase linearly until the second peak when the Ta layer yields and plastic deformation begins. The yield strength (the first peak), highest strength (the sec- ond peak) and mean flow strength of the three interface structures are shown in Table 2. It can be seen from the table that the KS{112} type sample has the highest flow strength, followed by the KS type sample, while the NW type sample has the lowest flow strength. Table 2: Yield strength, maximum strength and average flow strength of the three interfacial-structure samples Interface Loading orientation yield (GPa) max (GPa) flow (GPa) KS <112>Cu 7.75 7.85 4.22 NW <112>Cu 8.56 11.58 2.62 KS{112} <111>Cu 7.25 7.32 4.81 To explore the differences between the tensile me- chanical behaviors of the three samples of interface structures, this section studies the evolution law of mi- croscopic defects of the three samples during the tensile process. The microscopic defects at the yield point of Cu/Ta MNC Cu layers with three different interfacial structures are shown in Figure 4. To observe the change in the defect microstructure, the FCC and BCC structure atoms of the undeformed Cu layer and Ta layer are de- leted here. In all the samples, the Shockley partial dislo- cation (the green line in Figure 4, the Burgers vector is 1/6p) nucleation was observed on the Cu-layer side of the interface and in the form of partial dislocation loops along the {111} plane. The slip plane diffuses within the Cu layer, leaving an internal stacking fault (SF). In the KS-type Cu/Ta MNCs, before the Cu layer yields, part of the dissociated misfit dislocations diffuse in the Cu layer along the {111} slip plane parallel to the z-axis direc- tion, forming stacking faults penetrating the Cu layer. In the NW-type Cu/Ta MNCs, in addition to 1/6p, the for- mation of stepped rod dislocations (the red line in Fig- ure 4, the Burgers vector is 1/6n dislocation) was also observed. But for the KS{112}-type Cu/Ta MNCs, the balanced interface structure contains dissociative misfit dislocations, extending into the Cu layer, resulting in a lower activation barrier for dislocation nucleation at the interface. Therefore, at lower strain and stress, the for- mation of partial dislocation loops is first observed in the Cu layer. To release the stress and accommodate the strain, more Shockley partial dislocations nucleate and extend from the interface into the Cu layer. As the strain in- creases, the Shockley dislocation density increases, while the stress decreases for a short time. 22 The snap- shots of the microscopic defects of the three samples at the first peak-valley point are shown in Figure 5. Under this strain, part of the Shockley dislocation loop moves along the {111} slip plane to the other side of the inter- face, and the front segment of the dislocation loop anni- hilates, leaving only dislocation lines on both sides of the stacking fault. As the critical stress of the unit disloca- tion loop transport across the interface in the confine- ment layer model is higher than the critical stress of dis- location loop bending, and due to the difference between the geometric slip systems of the Cu and Ta layers, the dislocations at the interface cannot be transboundary and the planar dislocation transport is therefore blocked at the interface. According to Figure 5, sessile Lomer- Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... 418 Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 Figure 3: Stress-strain curves of Cu/Ta MNCs with different interface structures under tensile loading Figure 4: Snapshots of microscopic defects at the yield point of Cu layer: a) KS-type sample, = 0.065, b) NW-type sample, = 0.069, c) KS {112}, = 0.048 Cortell (LC) locks appear in the KS and NW-type sam- ples, while twinning is observed in the KS{112}-type sample. This indicates that the interfacial structure may affect the formation of dislocations and defects in the Cu layer, thereby affecting the way of plastic deformation. Due to the higher yield strength of Ta, dislocations in the Cu layer are blocked at the interface. Due to the weak interfacial shear, the shear stress field generated by the dislocations close to the interface traps the slip dislo- cations at the interface, so the dislocation density de- creases and the stress rises to the second peak point. Fig- ure 6 shows snapshots of the microscopic defects at the second stress peak (the yield point of the Ta layer). Ta layer yielding is divided into three stages: interfacial shear, accompanied by the movement of lattice disloca- tions in the interface plane; dislocation loop nucleation with concomitant interfacial structural compression; fi- nally, dislocations escape from the interface by overcom- ing the attraction within the interface layer. At the same time, it can be seen that the interface forms a 1/2o dislo- cation loop on the Ta-layer side. Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 419 Figure5: Snapshots of microscopic defects in the samples: a) KS-type sample, = 0.074, b) NW-type sample, = 0.077, c) KS{112}, = 0.058 Figure 7: Snapshots of defective microstructures during the formation of the Lomer-Cortell (LC) lock in the NW-type samples: a) = 0.068, b) = 0.069, c) = 0.070 Figure 6: Snapshots of microscopic defects at the yield point of the Ta layer: a) KS-type sample, = 0.086, b) NW-type sample, = 0.121, c) KS {112}, = 0.068 Figure 7 shows the formation of sessile Lomer- Cortell (LC) locks in the NW-type samples. 23,24 When the strain is 0.068, the interface forms a Shockley partial dis- location loop on the Cu-layer side (Figure 7a, the Bur- gers vector is 1/6p). As the leading portion slips, an in- trinsic stacking fault is formed in the slip region. Due to the intersection of the two dominant partial dislocations on the (111) plane, their motion is hindered and they merge to form ladder-rod dislocations. The merging pro- cess can be expressed with the following reaction for- mula: 1/6[2-1-1 ] +1/6[-121 ] 1/6[110 ] (1) The stepped rod dislocation is located at the intersec- tion of two different stacking faults where the Burgers vector is 1/6n and the slip plane is (100). Such disloca- tions cannot slide into the Cu layer and hinder the dislo- cation motion, thereby enhancing the mechanical proper- ties of the material. Two intersecting stacking faults form their Shockley partial dislocations and, in turn, sessile stepped rod dislocations at the intersections form sessile Lomer-Cortell (LC) locks. Figure 8 shows microscopic defect snapshots of the twin formation in the KS{112}-type samples. The unique interfacial structure of the KS{112}oCu||{112}nTa interface induced twinning due to the existence of a non-planar interface component and out-of-plane Burgers vector component of the interfacial dislocation. When = 0.048, the interfacial dislocations dissociate, and each dislocation decomposes into resid- ual dislocations and Shockley partial dislocations extend into the Cu layer. Subsequently, the Shockley dislocation partially slips along the {111} slip plane to form a stack- ing fault (SF 1 )( Figure 8a). A single layer of stacking faults cannot form twins, but it is the prerequisite for the formation of twins. The sliding of Shockley partial dislo- cations on adjacent planes creates twinning in FCC met- als. 25 So when the interval between SF 1 is small enough, the number of Shockley branches that subsequently need to be nucleated is reduced. On the contrary, when SF 1 is separated, the twins are more difficult to be excited, be- cause, from the energy point of view, it is difficult for the twins to nucleate in large numbers. When = 0.054, in the region adjacent to SF 1 , the second group of partial dislocations begins to nucleate and extend from the inter- face to the Cu layer (Figure 8b), resulting in a slippage on the {111} plane adjacent to the first fault, while the sliding surface is parallel to SF 1 . As the strain increases, another stacking fault (SF 2 ) forms near the first one; they are twinned when the strain in the region between SF 1 and SF 2 is 0.06 (Figure 8c). Unlike dislocation slip, de- formation twinning largely changes the lattice orienta- tion of the grains and forms adjacent boundaries between the twin orientation and the original lattice orientation. The twin interface acts as a substantial barrier to disloca- tion transport, preventing dislocation slippage. 26 There are out-of-plane dislocation defects and paral- lel defect arrays at the interface of the KS{112} type samples, so the nucleation dislocations tend to slip onto the parallel slip plane, and the Shockley partial disloca- tion motions are mostly of the first and second types. The interface of the KS-type and NW-type samples is a parallelogram or triangular defect array, and most of the nucleating dislocations move on different slip planes, but some Shockley dislocations tend to move more towards the third and fourth types. Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... 420 Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 Figure 9: Curves of dislocation density evolution with strain for the three interface structure samples Figure 8: Snapshots of defective microstructures of the twist formation in the KS {112} type samples: a) = 0.048, b) = 0.054, c) = 0.060 Figure 9 shows the evolution curves of dislocation density with strain in the samples with three interface structures. The dislocation density change is roughly di- vided into three stages according to the deformation pro- cess. The first stage is the elastic deformation stage. The dislocation density of the KS and NW samples is zero, while that of the KS{112} sample is not zero due to out-of-plane dislocations. Before the Cu layer of the sample yields, some out-of-plane dislocations are paral- lel to the z-axis. The expansion of the {111} plane and the breaking of out-of-plane dislocations lead to a de- crease in the dislocation density. The second stage is the yielding stage, where the yielding of the Cu layer leads to a rapid increase in the dislocation density, which leads to a decrease in the stress until the first peak and valley. During the first and second peaks of stress, the disloca- tion density decreases. This is because the dislocations close to the interface are trapped by the interface and some dislocations are annihilated due to the interactions of dislocations. After the Ta layer yields, the dislocation density increases rapidly. The third stage is the plastic deformation stage. The proliferation rate and disappear- ance rate of dislocations in the matrix tend to be in dy- namic equilibrium, and the dislocation density tends to be stable. 27 To reveal more details of the plastic deformation mechanism of the Cu/Ta MNCs with three interfacial structures, Figure 10 shows the corresponding radial dis- tribution functions before and after deformation. It can be seen that the peak intensities of the three curves be- fore and after deformation all decrease, and some long-range peaks disappear, indicating that the order of the samples decreases after deformation. For the KS and NW-type samples, the orderly decrease in the strength is attributed to the generation of dislocation and stacking fault defects, while for the KS{112}-type samples, the formation of twins is also a reason for the decrease in the strength. 28 4 CONCLUSION In this paper, the deformation mechanism of Cu/Ta nanolayered composites with different interface struc- tures was studied; the interaction between the interface structure and defect structure and the strengthening methods for the samples with different interface struc- tures were revealed. The main conclusion is that the KS{112} type sample interface is a parallel defect array, and the interface dislocation has a non-planar interface component, which induces deformation twinning. The formation of deformation twinning starts with the disso- ciation of intrinsic interfacial dislocations into Shockley partial dislocations, followed by the formation of SF 1 by dislocation motion; then, the second set of partial dislo- Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422 421 Figure 10: Radial distribution function curves before and after deformation of the samples with three interfacial structures cations may nucleate from the interface and slide on the adjacent SF 1 plane, and finally deformation twins are formed. The interface structures of the KS-type and NW-type samples are parallelogram and triangular inter- face defect arrays, respectively, which can easily induce two Shockley partial dislocations to slide along different (111) planes, intersect and form stepped rod dislocations. Acknowledgements This research was supported by the Guangdong Pro- vincial Key Field Research and Development Program (Grant No.: 2021B0707050001), and the scientific re- search team (Project No. HJL202202A007) of the Chaozhou Branch of the Guangdong Laboratory of Chemistry and Fine Chemicals (Hanjiang Laboratory). 5 REFERENCES 1 A. M. Abyzov, F. M. Shakhov, A. I. Averkin, V. I. Nikolaev, Mechan- ical properties of a diamond-copper composite with high thermal conductivity, Materials & Design, 87 (2015), 527–539, doi:10.1016/ j.matdes.2015.08.048 2 Z. L. Pan, Y. L. Li, Q. Wei, Tensile properties of nanocrystalline tan- talum from molecular dynamics simulations, Acta Materialia, 56 (2008) 14, 3470–3480, doi:10.1016/j.actamat.2008.03.025 3 Q. X. Pei, C. Lu, F. Z. Fang, H. Wu, Nanometric cutting of copper: A molecular dynamics study, Computational Materials Science, 37 (2006) 4, 434–441, doi:10.1016/j.commatsci.2005.10.006 4 J. H. Qin, Q. Chen, C. Y. Yang, Y. Huang, Research process on prop- erty and application of metal porous materials, Journal of Alloys and Compounds, 654 (2016) 39–44, doi:10.1016/j.jallcom.2015.09.148 5 K. A. Darling, E. L. Huskins, B. E. Schuster, Q. Wei, L. J. Kecskes, Mechanical properties of a high strength Cu-Ta composite at elevated temperature, Materials Science and Engineering A: Structural Mate- rials: Properties, Microstructure and Processing, 638 (2015), 322–328, doi:10.1016/j.msea.2015.04.069 6 A. S. Tran, T. H. Fang, Size effect and interfacial strength in nanolaminated Cu/CuxTa100-x composites using molecular dynam- ics, Computational Materials Science, 184 (2020), doi:10.1016/ j.commatsci.2020.109890 7 L. F. Zeng, R. Gao, Q. F. Fang, X. P. Wang, Z. M. Xie, S. Miao, T. Hao, T. Zhang, High strength and thermal stability of bulk Cu/Ta nanolamellar multilayers fabricated by cross accumulative roll bond- ing, Acta Materialia, 110 (2016), 341–351, doi:10.1016/j.actamat. 2016.03.034 8 G. P. Zhang, F. Liang, X. M. Luo, X. F. Zhu, A review on cyclic de- formation damage and fatigue fracture behavior of metallic nano- layered composites, Journal of Materials Research, 34 (2019)9 , 1479–1488, doi:10.1557/jmr.2019.22 9 X. Q. Liang, J. Y. Zhang, Y. Q. Wang, S. H. Wu, F. Zeng, K. Wu, G. Liu, G. J. Zhang, J. Sun, Tuning the size-dependent He-irradiated tol- erance and strengthening behavior of crystalline/amorphous Cu/Ta nanostructured multilayers, Materials Science and Engineering: A, 672 (2016), 153–160, doi:10.1016/j.msea.2016.07.005 10 X. Y. Zhu, J. T. Luo, F. Zeng, F. Pan, Microstructure and ultrahigh strength of nanoscale Cu/Nb multilayers, Thin Solid Films, 520 (2011) 2, 818–823, doi:10.1016/j.tsf.2010.12.251 11 H. Yavas, A. Fraile, T. Huminiuc, H. S. Sen, E. Frutos, T. Polcar, De- formation-Controlled Design of Metallic Nanocomposites, ACS Ap- plied Materials and Interfaces, 11 (2019) 49, 46296–46302, doi:10.1021/acsami.9b12235 12 T. Fu, X. Peng, S. Weng, Y. Zhao, F. Gao, L. Deng, Z. Wang, Molec- ular dynamics simulation of effects of twin interfaces on Cu/Ni multilayers, Materials Science and Engineering: A, 658 (2016), 1–7, doi:10.1016/j.msea.2016.01.055 13 S. Weng, H. Ning, N. Hu, C. Yan, T. Fu, X. Peng, S. Fu, J. Zhang, C. Xu, D. Sun, Y. Liu, L. Wu, Strengthening effects of twin interface in Cu/Ni multilayer thin films – A molecular dynamics study, Materials & Design, 111 (2016), 1–8, doi:10.1016/j.matdes.2016.08.069 14 R. Béjaud, J. Durinck, S. Brochard, Twin-interface interactions in nanostructured Cu/Ag: Molecular dynamics study, Acta Materialia, 144 (2018), 314–324, doi:10.1016/j.actamat.2017.10.036 15 C. Lv, J. Yang, X. P. Zhang, Y. Cai, X. Y. Liu, G. J. Wang, S. N. Luo, Interfacial effect on deformation and failure of Al/Cu nanolaminates under shear loading, Journal of Physics D: Applied Physics, 51 (2018) 33, doi:10.1088/1361-6463/aad2a8 16 S. Srinivasan, C. Kale, B. C. Hornbuckle, K. A. Darling, M. R. Chancey, E. Hernandez-Rivera, Y. Chen, T. R. Koenig, Y. Q. Wang, G. B. Thompson, K. N. Solanki, Radiation tolerance and micro- structural changes of nanocrystalline Cu-Ta alloy to high dose self-ion irradiation, Acta Materialia, 195 (2020), 621–630, doi:10.1016/j.actamat.2020.05.061 17 L. Lu, C. Huang, W. L. Pi, H. G. Xiang, F. S. Gao, T. Fu, X. H. Peng, Molecular dynamics simulation of effects of interface imperfections and modulation periods on Cu/Ta multilayers, Computational Mate- rials Science, 143 (2018), 63–70, doi:10.1016/j.commatsci.2017. 10.034 18 M. A. Bhatia, M. Rajagopalan, K. A. Darling, M. A. Tschopp, K. N. Solanki, The role of Ta on twinnability in nanocrystalline Cu-Ta al- loys, Materials Research Letters, 5 (2017) 1, 48–54, doi:10.1080/ 21663831.2016.1201160 19 R. F. Zhang, J. Wang, I. J. Beyerlein, T. C. Germann, Dislocation nu- cleation mechanisms from fcc/bcc incoherent interfaces, Scripta Materialia, 65 (2011) 11, 1022–1025, doi:10.1016/j.scriptamat.2011. 09.008 20 I. J. Beyerlein, J. Wang, K. Kang, S. J. Zheng, N. A. Mara, Twinn- ability of bimetal interfaces in nanostructured composites, Materials Research Letters, 1 (2013) 2, 89–95, doi:10.1080/21663831.2013. 782074 21 A. Stukowski, V. V. Bulatov, A. Arsenlis, Automated identification and indexing of dislocations in crystal interfaces, Modelling and Simulation in Materials Science and Engineering, 20 (2012)8 , doi:10.1088/0965-0393/20/8/085007 22 P. A. Gruber, E. Arzt, R. Spolenak, Brittle-to-ductile transition in ultrathin Ta/Cu film systems, Journal of Materials Research, 24 (2009) 6, 1906–1918, doi:10.1557/jmr.2009.0252 23 D. Choudhuri, A. CampBell, Interface dominated deformation mech- anisms in two-phase fcc/B2 nanostructures: Nishiyama-Wasserman vs. Kurdjumov-Sachs interfaces, Computational Materials Science, 177 (2020), doi:10.1016/j.commatsci.2020.109577 24 Z. G. Yan,Y. J. Lin, Lomer-Cottrell locks with multiple stair-rod dis- locations in a nanostructured Al alloy processed by severe plastic de- formation, Materials Science and Engineering A: Structural Mate- rials: Properties, Microstructure and Processing, 747 (2019), 177–184, doi:10.1016/j.msea.2019.01.066 25 M. Hosseini, N. Pardis, H. D. Manesh, M. Abbasi, D. I. Kim, Struc- tural characteristics of Cu/Ti bimetal composite produced by accu- mulative roll-bonding (ARB), Materials & Design, 113 (2017), 128–136, doi:10.1016/j.matdes.2016.09.094 26 C. B. Wang, J. J. Wang, J. N. Hu, S. L. Huang, Y. Sun, Y. L. Zhu, Q. Shen, G. Q. Luo, Shear Localization and Mechanical Properties of Cu/Ta Metallic Nanolayered Composites: A Molecular Dynamics Study, Metals, 12 (2022) 3, doi:10.3390/met12030421 27 B. B. He, B. Hu, H. W. Yen, G. J. Cheng, Z. K. Wang, H. W. Luo, M. X. Huang, High dislocation density-induced large ductility in de- formed and partitioned steels, Science, 357 (2017) 6355, 1029–1032, doi:10.1126/science.aan0177 28 L. Lu, Y. F. Shen, X. H. Chen, L. H. Qian, K. Lu, Ultrahigh strength and high electrical conductivity in copper, Science, 304 (2004) 5669, 422–426, doi:10.1126/science.1092905 Z. HAOTIAN et al.: MOLECULAR DYNAMICS SIMULATION OF TENSILE DEFORMATION OF ... 422 Materiali in tehnologije / Materials and technology 56 (2022) 4, 415–422