# Fileset

[nanomaterials-15-01528 (1).pdf](https://mdr.nims.go.jp/filesets/b3ab5063-f4bb-4cb6-a00f-9f337eb5249f/download)

## Creator

[Rahul Sahay](https://orcid.org/0000-0002-0019-0153), Ihor Radchenko, Pavithra Ananthasubramanian, Christian Harito, [Fabien Briffod](https://orcid.org/0000-0002-3635-4885), Koki Yasuda, Takayuki Shiraiwa, Mark Jhon, Rachel Speaks, [Derrick Speaks](https://orcid.org/0000-0003-3540-1720), Kangjae Lee, Manabu Enoki, Nagarajan Raghavan, Arief Suriadi Budiman

## Rights

[Creative Commons BY Attribution 4.0 International](https://creativecommons.org/licenses/by/4.0/)

## Other metadata

[Interface Rotation in Accumulative Rolling Bonding (ARB) Cu/Nb Nanolaminates Under Constrained and Unconstrained Loading Conditions as Revealed by In Situ Micromechanical Testing](https://mdr.nims.go.jp/datasets/344651e5-33ef-4803-a468-d3216b85f27c)

## Fulltext

Interface Rotation in Accumulative Rolling Bonding (ARB) Cu/Nb Nanolaminates Under Constrained and Unconstrained Loading Conditions as Revealed by In Situ Micromechanical TestingAcademic Editor: Julian MariaGonzalez EstevezReceived: 30 June 2025Revised: 15 August 2025Accepted: 26 August 2025Published: 7 October 2025Citation: Sahay, R.; Radchenko, I.;Ananthasubramanian, P.; Harito, C.;Briffod, F.; Yasuda, K.; Shiraiwa, T.;Jhon, M.; Speaks, R.; Speaks, D.; et al.Interface Rotation in AccumulativeRolling Bonding (ARB) Cu/NbNanolaminates Under Constrainedand Unconstrained LoadingConditions as Revealed by In SituMicromechanical Testing.Nanomaterials 2025, 15, 1528. https://doi.org/10.3390/nano15191528Copyright: © 2025 by the authors.Licensee MDPI, Basel, Switzerland.This article is an open access articledistributed under the terms andconditions of the Creative CommonsAttribution (CC BY) license(https://creativecommons.org/licenses/by/4.0/).ArticleInterface Rotation in Accumulative Rolling Bonding (ARB)Cu/Nb Nanolaminates Under Constrained and UnconstrainedLoading Conditions as Revealed by In SituMicromechanical TestingRahul Sahay 1,2 , Ihor Radchenko 1,3, Pavithra Ananthasubramanian 2, Christian Harito 4, Fabien Briffod 5,Koki Yasuda 6, Takayuki Shiraiwa 6, Mark Jhon 7, Rachel Speaks 8,9, Derrick Speaks 8,9 , Kangjae Lee 9,Manabu Enoki 6, Nagarajan Raghavan 2 and Arief Suriadi Budiman 4,8,9,*1 Xtreme Mechanics Laboratory, Engineering Product Development (EPD), Singapore University of Technologyand Design (SUTD), 8 Somapah Road, Singapore 487372, Singapore2 Nano-Micro Reliability Laboratory (MRL), Engineering Product Development, Singapore University ofTechnology and Design (SUTD), Singapore 487372, Singapore3 Center for Advancing Materials Performance from the Nanoscale (CAMP-Nano), State Key Laboratory forMechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, China4 Industrial Engineering Department, BINUS Graduate Program—Master of Industrial Engineering, BinaNusantara University, Jakarta 11480, Indonesia5 Research Center for Structural Materials, National Institute for Materials Science, 1-2-1 Sengen,Tsukuba 305-0047, Ibaraki, Japan6 Department of Materials Engineering, School of Engineering, The University of Tokyo, 7-3-1 Hongo,Bunkyo-ku, Tokyo 113-8656, Japan; enoki@rme.mm.t.u-tokyo.ac.jp (M.E.)7 Institute of High Performance Computing (IHPC), Agency for Science, Technology and Research (A*STAR), 1Fusionopolis Way, #16-16 Connexis, Singapore 138632, Singapore8 Department of Manufacturing and Mechanical Engineering and Technology (MMET), Oregon Institute ofTechnology, Klamath Falls, OR 97601, USA9 Oregon Renewable Energy Center (OREC), Oregon Institute of Technology, Klamath Falls, OR 97601, USA* Correspondence: suriadi@alumni.stanford.eduAbstractAccumulative rolling bonding (ARB) Cu/Nb nanolaminates have been widely observedto exhibit unique and large numbers of interface-based plasticity mechanisms, and thesehave been associated with the many extraordinary properties of the material system,especially resistances in extreme engineering environments (mechanical/pressure, thermal,irradiation, etc.) and ability to self-heal defects (microstructural, as well as radiation-induced). Recently, anisotropy in the interface shearing mechanisms in the material systemhas been observed and much discussed. The Cu/Nb nanolaminates appear to shear onthe interface planes to a much larger extent in the transverse direction (TD) than in therolling direction (RD). Related to that, in this present study we observe interface rotationin Cu/Nb ARB nanolaminates under constrained and unconstrained loading conditions.Although the primary driving force for interface shearing was expected only in the RD,additional shearing in the TD was observed. This is significant as it represents an interfacerotation, while there was no external rotational driving force. First, we observed interfacerotation in in situ rectangular micropillar compression experiments, where the interface issimply sheared in one particular direction only, i.e., in the RD. This is rather unexpected as,in rectangular micropillar compression, there is no possibility of extra shearing or drivingforce in the perpendicular direction due to the loading conditions. This motivated us tosubsequently perform in situ microbeam bending experiments (microbeam with a pre-made notch) to verify if similar interface rotation could also be observed in other loadingmodes. In the beam bending mode, the notch area was primarily under tensile stress in theNanomaterials 2025, 15, 1528 https://doi.org/10.3390/nano15191528https://doi.org/10.3390/nano15191528https://doi.org/10.3390/nano15191528https://creativecommons.org/licenses/by/4.0/https://creativecommons.org/licenses/by/4.0/https://www.mdpi.com/journal/nanomaterialshttps://www.mdpi.comhttps://orcid.org/0000-0002-0019-0153https://orcid.org/0000-0003-3540-1720https://doi.org/10.3390/nano15191528https://www.mdpi.com/article/10.3390/nano15191528?type=check_update&version=1Nanomaterials 2025, 15, 1528 2 of 21direction of the beam longitudinal axis, with interfacial shear also in the same direction.Hence, we expect interface shearing only in that direction. We then found that interfacerotation was also evident and repeatable under certain circumstances, such as under anoffset loading. As this behaviour was consistently observed under two distinct loadingmodes, we propose that it is an intrinsic characteristic of Cu/Nb interfaces (or FCC/BCCinterfaces with specific orientation relationships). This interface rotation represents anotherinterface-based or interface-mediated plasticity mechanism at the nanoscale with importantpotential implications especially for design of metallic thin films with extreme stretchabilityand other emerging applications.Keywords: multilayers; nanolaminates; interface-based plasticity mechanism; nanoplasticity1. IntroductionBimetallic nanolaminates have been experimentally analyzed, simulated, and mathe-matically modeled to fulfil applications for strong/tough, wear-resistant, fracture-resistant,impact-resistant, and advanced structural materials [1,2]. A large number of interfaces innanolaminates have the potential to modify bulk/local interfacial plasticity and thereby itsstrength and toughness as well as fracture/impact resistance [3,4]. Here, the test materialis ARB Cu/Nb nanolaminate, which is being extensively considered for its remarkableflow strength, deformability, thermal stability, and radiation shielding due to its uniquesemi-coherent immiscible interface [5–7]. Cu/Nb nanolaminate has lattice mismatch whichprovides the capability to stop, absorb, and annihilate dislocation at its interface to mitigatestrain localization under applied loading [8,9]. Shao et al. [10] noted that the existenceof residual compressive stresses in Nb and residual tensile stresses in Cu may make dis-locations bow in opposite directions in Cu/Nb nanolaminates. Zhang et al. [11] furtherhighlighted the role of the interface in the plastic behavior of Cu/Nb nanolaminates undershock compression testing mode. They noted that critical shock pressures for nucleatingand transmitting dislocations through a flat interface are considerably higher than thosefrom a faceted interface owing to interface characteristics causing significantly diversemechanisms for dislocation nucleation, absorption, and transmission.Furthermore, the Cu/Nb ARB nanolaminate displays crystallographic anisotropyowing to the distinct interfaces along the rolling direction (RD) and the transverse direction(TD) [12,13]. Molecular dynamics simulations predicted that the Cu/Nb ARB nanolam-inate’s interface is highly anisotropic and theoretically has highly prohibitive interfacialshear in the rolling direction and has a finite interface shear strength ~1.2 GPa in thetransverse direction (TD) [14]. Budiman et al. [15,16] highlighted the role of interfaces inhindering crack propagation and affecting plasticity mechanisms. Radchenko et al. [16]further emphasize the impact of interface shear strength on fracture behavior, with dif-ferences observed amongst the transverse and rolling directions. Along the transversedirection (TD) of Cu/Nb ARB nanolaminate, Budiman et al. [15–18] noted notch wideningduring cross-layer fracture, indicating a significant extent of interfacial sliding/shearingattributed to size effects and localized shearing/plasticity mechanisms on the interfaces.Radchenko et al. [16] documented improved plasticity with low interface shear strengthalong the TD compared to the RD.Local plastic instability due to strain localization has been observed experimentallyduring Cu/Nb nanolaminate micropillar compression. Zheng et al. [19] showed that, withinthe shear band formed during micropillar compression, the layers in Cu/Nb nanolaminatecould rotate to orientations favorable for interfacial sliding. This phenomenon has alsoNanomaterials 2025, 15, 1528 3 of 21been observed in Cu/Nb ARB nanolaminates, indicating the significance of interfacesin the enhancement the ductility of such nanolaminated materials [12]. Furthermore,Demkowicz et al. [14] predicted the possibility of rotation of interfaces through moleculardynamic simulations during micropillar compression tests.As the occurrence and propensity of interfacial plasticity along the TD is well ob-served [14,15,20,21], the intention of the current study is to document interfacial shearingalong the RD in Cu/Nb ARB nanolayered samples, as well as to report if additional inter-facial plasticity mechanisms such as rotation are evident and under which circumstances.Once the conditions under which interface shear/rotation along the RD are well docu-mented, the insight could support the designer/researcher in their quest for designingstructures or components of Cu/Nb nanolaminated materials workable under varyingloading configurations/conditions. To achieve the above outcome, two sets of experi-ments were designed: (1) in situ rectangular micropillar compression essentially to achieveunconstrained interfacial shear/rotation and study its impact on interface-based local-ized shearing/plasticity mechanisms along the RD and (2) in situ microbeam bendingtests of pre-notched Cu/Nb (63 nm) ARB clamped beams to examine constrained local-ized interfacial plasticity along the TD and interfacial rotation along the RD under offsetloading. In the latter study, offset loading was introduced to examine the effect of localbending moment/bending load on localized interfacial plasticity (shear/rotation). Thestudy documents Cu/Nb nanolaminate’s localized interface-based plasticity mechanismsduring in situ rectangular micropillar compression and in situ microbeam bending, essen-tial for interface engineering to enhance the ductility and achieve extreme stretchabilityof nanolayered composites. These outcomes of the study could be very significant foradvanced structural applications, as they offer understandings of applied loading config-urations/conditions and their effect on the interfacial shear/rotation along the TD andRD for anisotropic Cu/Nb nanolaminate. The authors think that information regardinginterfacial sliding/rotation under diverse loading configurations could help researchersin scheming workable design diagrams for a given material system, which could moti-vate the design/fabrication of novel strong/tough assemblies sustainable under complexloading configurations/conditions for evolving functionalities, like stretchable bimetallicconductors for innovative wearable devices.2. Materials and Methods2.1. Assembly of ARB Cu/Nb NanolaminatesCu/Nb nanolaminates are created through accumulative roll bonding where alternat-ing layers of pre-treated Cu and Nb are stacked and later rolled together. The rolled layersare then trimmed and stacked again and rolled repeatedly in the same rolling direction tillthe distinct layer thickness is shortened down to nanometers (63 nm for the test material inthe work) [20]. During the ARB, the plastic deformation is performed along the ND. Asthe volume of the sample is conserved, the sample suffers expansion (rolling reductionof 99.999%) in the RD and no deformation along the TD. Therefore, the ARB producesdifferent textures along rolling and transverse directions and thereby distinct interfacialcharacteristics [8,21–24]. Additional specifics of the accumulative roll bonding and theCu/Nb texture development can be found in the literature [2,25].2.2. Fabrication of Cu/Nb ARB Rectangular Micropillar SamplesThe Cu/Nb ARB nanolaminate rectangular micropillars for in situ pillar compres-sion experiments were fabricated using FIB milling (FEI Nova Nanolab dual-beam FIB,Hillsboro, OR, USA) at Nanyang Technological University’s Microelectronics Reliability &Characterization Laboratory in Singapore. The milling process applied for the fabricationNanomaterials 2025, 15, 1528 4 of 21of micropillars consists of three milling stages, (a) a constant 30 kV acceleration voltagewith 21 nA beam current for the first two stages of the milling and (b) in the third stage,a constant 30 kV acceleration voltage was applied, whereas the beam current was gradu-ally reduced to 28 pA from 21 nA for polishing the sample. Two pillars were made withinterfaces inclined 45◦ with regard to the loading axis. In this study, micropillars withrectangular cross-sections were used in contrast to the more typical circular pillars reportedthus far in the literature [4,26–30]. The rectangular cross-section allows the direct testing ofinterface shear strength along the RD as well as the TD of the Cu/Nb ARB nanolaminate,as illustrated in Figure 1a and shown in Figure 1b,c. The width to height ratios (1:2 and1:3) were selected to suppress buckling instability during pillar compression, as suggestedby Zhang et al. [31]. Similar aspect ratios (1:2 and 1:3) were also used by Li et al. in theirinterface shear study of Cu/Nb PVD nanolaminates [32]. Figure 1. (a) Schematic of the in situ rectangular pillar compression experiments. SEM images ofas-fabricated pillars with ±45◦ offset from (b) RD and (c) TD, respectively.2.3. Fabrication of Cu/Nb ARB Microbeam SamplesIn order to fabricate Cu/Nb ARB microbeam samples, FEI Nova Nanolab dual-beamFIB at Nanyang Technological University’s (NTU) Microelectronics Reliability & Characteri-zation Laboratory, Singapore was employed to manufacture clamped microbeams. FIB wasused to prepare microbeams from pre-polished nanolaminates with ~63 nm layer thickness(see Figure 2). Initially, Cu/Nb ARB was milled at 30 kV with a high beam current toproduce a clamped microbeam configuration. Subsequently, a low milling current was usedfor the finer cutting and polishing to the final dimensions. The center notches were milledat 10 pA current as a single line cut across the width of the beam. The notch depths werekept between 1/5 and 1/3 of the beam heights, whereas the notch widths ranged between100 nm and 360 nm. Figure 2 shows representative images of the fabricated clampedmicrobeam. The microbeam dimensions were selected to satisfy both small-scale yieldingNanomaterials 2025, 15, 1528 5 of 21and plane strain [33]. The small gradient height (height/thickness) typically permits stablecrack expansion, whereas the large gradient height may prevent edge cracking for theabove clamped microbeam configuration. Figure 2. SEM images depicting front view of Cu/Nb (TD) Beam, (a) the full-frontal view of thebeam with dimensions: L~40 µm, W~5 µm, and T~5 µm, (b) close-up of the frontal notch in the beamfabricated by FIB with dimensions: H = 1240 nm, W = 120 nm.2.4. In Situ Microbeam BendingThe in situ microbeam bending tests were performed by means of a PI-85 SEM pico-indenter (Bruker®, Billerica, MA, USA) in JEOL FESEM (JSM-7600F), Tokyo, Japan. Themicrobeams were loaded by a truncated cone-shaped indenter with a top surface withdiameter 5 µm. The bending load was applied using a displacement control mode througha fixed 5 nms−1 displacement rate. The generated load–displacement data was correctedfor thermal drift, which was estimated to be of the order of ~10 µN. Two sets of experimentswere performed: (a) microbeam bending with no offset loading (Cu/Nb-NOL), wherethe indenter tip is aligned along the notch of the microbeam and (b) microbeam bendingwith offset loading (Cu/Nb-OL), where the indenter tip loads the sample with an offsetfrom the center of the notch along the length of the beam to generate a combination oflocal bending moment at the notch tip and bending load at the notch of the microbeam.The load–displacement data were thus generated and are discussed in the subsequentsection below.For both in situ micropillar compression and in situ microbeam bending tests, theload–displacement curve and corresponding real-time video were synchronized and takenusing a frame grabber with the TriboScan software version 9 (for TI-950) (Bruker®, USA).Also, before the experiments, drift correction measurements were performed to then correctthe final load–displacement measurements.3. Experimental Results3.1. In Situ Rectangular Micropillar CompressionThe crystallographic anisotropy in Cu/Nb nanolaminates leads to different interfacesalong the RD and TD [12,13]. Typically, in situ rectangular micropillar compression experi-ments provide a well-defined stress state, making them suitable for studying anisotropyeffects [34]. Additionally, apart from mechanical anisotropy, in situ micropillar compres-sion could also provide more direct insights into mechanical behavior earlier and/or later,surpassing the elastic limit until rupture [35]. Furthermore, the response of the materialNanomaterials 2025, 15, 1528 6 of 21during micropillar compression tests can be quantitatively mapped to understand the stressand strain distribution at small length scales [36].In the current work on Cu/Nb ARB nanolaminate samples, in situ rectangular mi-cropillar compression experiments were performed along the RD as well as in TD to probethe interfacial plasticity mechanism. Molecular dynamics simulations have predicted thatthe Cu/Nb ARB interface is highly anisotropic and theoretically does not shear in the RDand has a finite interface shear strength of ~1.2 GPa in the TD [14].The results of in situ rectangular micropillar compression experiments along the TDand RD are shown in Figure 3. Their corresponding engineering-resolved shear stress vs.displacement plots (see Figure 4). The engineering-resolved shear stress vs. displacementcurves were attained from the experimentally obtained load vs. displacement curves usingσeng = −F(h0 − ∆)/A0h0 engineering-resolved shear stress if the shear direction followsthe RD or TD according to the interface inclination, where σeng is engineering stress, F isthe measured load, ∆ is the measured displacement, A0 is an initial cross-section area of thepillar, and h0 is the initial pillar height.Figure 3. SEM images of the deformed TD (a,b) and RD (c,d) rectangular micropillars. The TDmicropillar exhibits the interface shear along two interfaces of a single Cu layer, as can be seen fromthe back-scattered electron image in (a). The shear does not result in free Nb surface creation. Onlythe Cu layer surface is exposed, as can be seen from (b). In the case of the RD micropillar, the toppart of the RD micropillar is rotated relative to its bottom part, as shown in (d). One layer near therotation region is broken, as shown in the inset in (c). The micropillar was sheared in RD, and due toextreme stress concentration built up at one of the corners of the pillar, the shearing in the RD wasresisted, and then the whole pillar was stuck in that corner, thus activating deformation of materialsin the other direction (i.e., TD), thus the interface rotates.Nanomaterials 2025, 15, 1528 7 of 21Figure 4. Engineering-resolved shear stress along the interface vs. indenter displacement for theTD and RD rectangular micropillars as shown in Figure 3: (a) The interface shear in the TD pillaris activated at ≈300 MPa shear stress followed by stress increase with deformation. The stress isunderestimated here due to the decrease in the contact area upon pillar deformation and (b) theinterfacial rotation (evident from Figure 3c,d) in the RD pillar occurs at ≈155 MPa immediatelyfollowed by decrease in stress.3.1.1. Interface Shear in the TD Rectangular Micropillar CompressionThe ARB 63 nm Cu/Nb TD micropillar exhibits interfacial shear, which starts to occurat ≈300 MPa (see Figure 4a), which is much smaller than the 1000–1200 MPa predictedvia MD simulations [14]. This may be attributed to its finite interface in the experimentcompared to infinite interface modeled in the simulation. It is observed that dislocationsmoving out from a finite interface would encounter fewer misfit dislocations compared toan infinite interface modeled through MD simulation. The shear was only observed at oneside of the interface, where a free Cu surface was formed after the shear (see Figure 3a,b).The inhomogeneity of the interface shear implies that the interface shear strength dependson the boundary conditions, and it is lower when a Cu free surface is formed and higherwhen a Nb free surface is formed.The interface shear observed at two of the interfaces is initiated at ≈300 MPa and thestress keeps increasing with further deformation (see Figure 3a). Each load drop in the loadvs. displacement curves corresponds to dislocation nucleation at the interface trailed bythe dislocations leaving the interface [37,38]. The plasticity in the Cu layer is governed bythe CLS and leads to plastic hardening. This plastic hardening is likely to be attributed tothe observed increase in the stress (see Figure 4a). Further, the Cu plasticity can result inadditional dislocations deposited at the shearing interface. Dislocations deposited at theCu/Nb interface can lead to dislocation dissociation, which can affect the dislocation glidealong the interface which can also contribute to the observed increase in the shear stress.3.1.2. Interface Plasticity in the RD Rectangular Micropillar CompressionIn the case of the 63 nm Cu/Nb micropillar, the shear along the RD is theoreticallyprohibited due to extreme interfacial shear strength in this direction. Nevertheless, theinterfacial shear along the RD is observed in Figure 3c,d. Later, because of high stressconcentration built up at one of the corners of the micropillar, the shearing in the RD wasresisted, thus activating the interface shearing along the TD, and hence the interface rotationobservation. The nanolayer rotation towards the TD can be associated with the creation ofstacking faults, which are perpendicular to the interface plane. The stacking fault is formedNanomaterials 2025, 15, 1528 8 of 21as a result of an extension of a partial dislocation, which originated from misfit dislocationswith the Burgers vector perpendicular to the Cu/Nb ARB interface [14,39]. However, therotation occurs at ≈155 MPa, which is below the theoretical shear stress of the stackingfault extension (≈300 MPa) [14,39]. This low critical stress can be attributed to pre-existingstacking faults in the Cu/Nb ARB rectangular micropillar. These pre-existing stacking faultsallow a new stacking fault formation at low shear stress, as proposed by Zheng et al. [40].The engineering shear stress vs. displacement curve for the RD micropillar shows that theshear stress is reduced to ≈100 MPa after the initiation of the layer rotational deformation(see Figure 3b). The major engineering stress drop in the RD micropillar can be associatedwith Cu layer failure, as shown in the inset in Figure 3c. Further deformation corresponds toa reduction of the contact area between the top and bottom part of the RD micropillar, whichmeans that true shear stress is larger than the calculated ≈100 MPa engineering-resolvedshear stress.As shown in Figure 4, a comparison of the engineering-resolved shear stress versusindenter displacement for the TD and RD rectangular micropillars indicates that the RDmicropillar exhibits yielding at a significantly lower stress (~155 MPa) compared to the TDmicropillar (~300 MPa). As has been discussed more completely in the previous paragraph,it was due to high stress concentration built up at one of the corners of the micropillaras the shearing in the RD was resisted, thus activating the interface shearing along theTD, and hence the interface rotation observation. Hence, it should not be interpreted thatshearing in the RD is easier than shearing in the TD. According to the molecular dynamicsimulations by Wang et al. [41] and Demkowicz et al. [14], the interface shear strength alongthe RD is almost infinite for ARB Cu/Nb nanolaminate, though in the TD, interface shearstrength would have a finite value. The RD micropillar was in effect plastically deformedat ~155 MPa via local plastic instability mechanisms other than the pure interface shearingin the rolling direction (RD) that the experiment was designed for.These observations above were representative of a few rectangular micropillar com-pression experiments (4–5 micropillar samples) that we had conducted for each TD andRD micropillar compression in the present study. We were limited by the prohibitivelyhigh cost of the rectangular micropillar sample fabrication (using a focused ion beam (FIB)).Nevertheless, within this constraint, we observed the above phenomena (in TD as well asin RD micropillar compressions) consistently. All the TD micropillar compression experi-ments did not exhibit interfacial rotation—just straightforward interfacial shearing leadingto failure (the shear yielding points may differ slightly—within a range of ±~100 MPa).Meanwhile, all the RD micropillar compressions ended up with interfacial rotation—againextent and actual rotating nanolayers may differ slightly from sample to sample, as well asin terms of the shear stress–displacement curves (qualitative shape, as shown in Figure 4b)and yield points (they differ within ±~150 MPa) in the present study. This could be as-sociated with the pre-existing defects and other actual sample to sample variations, ashas been described above. Interfacial rotation is clearly an important part of the overallinterface-based plasticity mechanisms in the Cu/Nb ARB nanolaminate samples.3.2. In Situ Cu/Nb ARB Microbeam BendingIn situ microbeam bending experiments were performed on Cu/Nb ARB nanolam-inates (with layer thickness ~63 nm) to probe localized interfacial plasticity and frac-ture/failure resistance via crack growth/propagation through the thickness of the sample.It has been well documented that the layer thickness controls bulk and interfacial plasticityin bimetallic nanolaminates [22]. For ∼10 < h < ∼100 nm, where h is the distinct layerthickness, the motion of dislocation is typically governed by confined layer slip (CLS),which encompasses the motion of sole dislocation loops parallel to the interfaces withinNanomaterials 2025, 15, 1528 9 of 21layers [24,42,43]. In the present work, as the layer thickness being studied is 63 nm, dislo-cation mechanisms may be expected to be governed by confined layer slip (CLS) whichcould induce enhanced ductility/toughness owing to limited plastic strain localization [44].Apart from bulk localized plasticity of the constituent material layers, localized interfacialplasticity contributes to the overall plasticity of the nanolaminates. Interfacial plasticitywhich happens via nucleation and glide of interface dislocations parallel to the interfaceplane contributes to the overall plasticity [45]. In addition, in situ microbeam bendingexperiments were executed on the Cu/Nb ARB samples clamped along the TD—thatis, to study the local interfacial plasticity mechanisms (near the notch) while the overallmicrobeam samples were stretched (pulled, in tension) along the transverse direction (TD)of the Cu/Nb ARB samples. From previous work [12,13], we may expect that interfacialshearing or sliding here would be much less inhibited by the dislocation configurations onthe interfaces, compared to along the RD of the Cu/Nb ARB samples. This is related tothe much higher interfacial shear strength along the RD (compared to TD), as predicted byDemkowicz et al. with molecular dynamics simulations [14].3.2.1. In Situ Microbeam Bending of Cu/Nb ARB Microbeam with No Offset Loading(Cu/Nb-NOL)In situ microbeam bending of the Cu/Nb TD microbeam with no offset loading(Cu/Nb-NOL, Figure 5) exhibits notch widening, followed by crack initiation and prop-agation to final failure [12,13]. Along the TD, the Cu/Nb-NOL has low interfacial shearstrength. The low interfacial shear stress along the TD can shear the interface in responseto the stress fields generated by an adjoining lattice dislocation and could induce thedislocation in the interface [9,46]. These absorbed dislocations (irrespective of their sign)could extend their core inside the interface plane, which could result in localized interfacialplasticity. Also, these absorbed dislocations require re-nucleation in order to glide on theoutward-bound slip system in the adjacent nanolayer. Therefore, the Cu/Nb-NOL TDmicrobeam provides stronger resistance to dislocation transmission and high localizedinterfacial plasticity and resistance to crack propagation [9,23].The degree of notch widening for Cu/Nb-NOL was ~900% with respect to the initialnotch width. Notch widening was calculated based on the original notch width (120 nm),and the final notch before the crack starts to propagate out of the notch corner across thethickness of the Cu/Nb-NOL. Typically, when the dislocations generated from the exter-nally applied loading get deposited in the Cu/Nb interface, dislocation-mediated plasticitymechanisms get activated within the interface, which results in localized interfacial shear.Hence, the degree of notch widening discussed could be associated with interfacial shearstrength/interfacial shear of the Cu/Nb nanolaminate.In the case of Cu/Nb-NOL TD microbeam bending, high localized interfacial plasticityis typically observed [9,23]. Normally, in the instance of the Cu/Nb-NOL TD beam withoutany designed offset loading (in this work) or torque (in our previous work [47]), complexstress states do not develop. In general, the interface shear near the notch tip mostly followsthe nominal stress state due to beam bending, i.e., layers of the Cu/Nb-NOL TD slide overeach other near the tip of the notch, and this shearing action follows the general stressdistribution caused by the bending of the beam. This is consistent with what we observedin this experiment (Cu/Nb-NOL). No significant interfacial rotation was observed near thelocation of the tip notch.As the load is further increased, the dislocation starts to pile up at the corner of thenotch to initiate crack. The propagation of crack through the thickness of the nanolam-inate requires continuous generation and accumulation of dislocations. Nevertheless,dislocations may get absorbed, deposited, and annihilated in the lattice mismatchedNanomaterials 2025, 15, 1528 10 of 21Cu/Nb-NOL TD interface, resulting in interfacial shear via dislocation-mediated local-ized plasticity mechanisms, thus hindering the propagation of the crack through thethickness of the microbeam. Subsequently, with an increase in the applied load, the crackwould propagate, resulting in shear band formation and final failure. Nevertheless, inthis particular case, crack propagation and subsequent shear band formation were hin-dered due to the failure of the microbeam at one of its fixed ends. Similar effects of lowinterfacial shear strength along the TD of Cu/Nb-NOL leading to localized interfacialplasticity and hindrances to the propagation of the crack, as well as the no significantinterfacial rotation observation, have been previously reported [48,49] and establishedby means of finite element analysis [16]. Figure 5. Load–displacement curve of ARB 63 nm Cu/Nb microbeam bending test with no offset load-ing (NOL) along TD: (1) the notch widens up until the instigation of crack, (2) additional notch widen-ing with slight crack growth, (3) crack growth with notch widening, and (4) crack growth and forma-tion of shear instabilities were suppressed due to the failure at one of the beam’s edges. No significantinterfacial rotation was evident. Similar observations have been previously reported [15,48,49].As can be expected, during in situ microbeam bending, the interactions betweendislocations generated due to the external load and interfaces could adjust the interface con-struction and orientation of slip systems across the interface, which could then potentiallylead to some degree of interfacial rotation under certain circumstances. At substantiallyhigher externally applied loads, these interfacial rotations do occur, which are essentialfor the alignment of slip systems for the flow of dislocations across the layers of Cu/Nbnanolaminate. Therefore, to a certain extent, the interfacial rotation is always possible in insitu microbeam bending with no offset loading or torque (Cu/Nb-NOL). Nevertheless, as isdiscussed subsequently in the next section (Section 3.1.2 Cu/Nb OL), the offset loading waspurposely introduced in the instance of the Cu/Nb-OL microbeam (with offset loading) togenerate complex stress states and hence subsequently leading to interfacial rotation.Nanomaterials 2025, 15, 1528 11 of 213.2.2. In Situ Microbeam Bending Experiments of Cu/Nb ARB Microbeam withOffset Loading (Cu/Nb-OL)It has been noted in Section 3.2.1 that Cu/Nb-NOL showed appreciable localizedinterface plasticity along the TD, therefore, the aim here is to examine the significance of anoffset loading on interfacial rotation, localized interface plasticity, and crack initiation andpropagation in Cu/Nb nanolaminate (see Figure 6). Figure 6. In situ microbeam bending tests were executed on pre-notched 63 nm Cu/Nb ARB clampedbeam along the transverse direction (TD) with an offset loading. SEM image depicts a full beam withdimensions: Length = 40 µm, Width = 5.0 µm, and Thickness = 4.8 µm. The image depicts an offsetloading applied to the beam with an Offset = 0.1 * Length (4 µm). The hypothetical dotted load linesare plotted to show that the offset loading results in a normal load at the notch tip (responsible for thebending of the beam) and a local bending moment restricting the flow of dislocations along the TD.The introduction of offset loading induces the mixed mode loading in Cu/Nb-OL,which is otherwise similar to Cu/Nb-NOL in both in situ beam bending experimentalparameters as well as the nominal beam geometry. Apart from the normal load componentto induce the shear along the monolayers in the TD similar to Cu/Nb-NOL, a local bendingmoment is also introduced. The local bending moment was introduced to restrict the shearalong the monolayers and thereby interfacial/bulk plasticity along the TD. The aim was tostudy if restriction of interfacial/bulk plasticity along the TD may induce interfacial shearalong the RD and hence interfacial rotation.Cu/Nb-OL (Cu/Nb microbeam with offset loading) goes through similar regimesof notch widening and crack initiation, followed by its propagation and subsequent finalfailure (see Figures 5 and 7). Nevertheless, the degree of the notch widening was rather low,around ~400%, compared to Cu/Nb-ARB-NOL (around ~900%). It is acknowledged thatthe motion of dislocations and thereby dislocation-mediated plasticity have a significantpart in the deformation behavior of these Cu/Nb nanolaminates. Typically, dislocationtransmission could occur across or along the interface of Cu/Nb nanolaminates, whichcould lead to the formation of shear band or interfacial shear, respectively, while in the casewhere dislocations do not transmit across or along the interface, interface tilting or rotationNanomaterials 2025, 15, 1528 12 of 21could occur [19]. In the present case, due to incorporation of the local bending moment atthe notch tip, the interfacial shear was relatively more constrained along the TD (comparedto Cu/Nb NOL), thus promoting interfacial shear along the RD (see Figure 8) and hencelocal interfacial rotation near the notch tip was evident (see Figure 9).Figure 7. In situ microbeam bending test of 63 nm Cu/Nb TD microbeam with offset loading (OL):(a,b) the original notch widening, (c) notch widening trailed by crack instigation, (d–f) additionalnotch widening with crack evolution and spread through several layers from the notch tip, followedby eventual failure of the beam. Corresponding load–displacement curve is shown in Figure 8. Localinterfacial rotation near the notch tip was evident (larger magnification image is provided in Figure 9).Figure 8. Load–displacement plot of 63 nm Cu/Nb TD microbeam (Cu/Nb-OL). The load riseswith the applied displacement of the indenter tip due to the applied loading configuration shown inFigure 6, which constrains the interfacial shear along the TD but promotes the rotation along RD andsubsequently shear band formation and eventual failure of the beam.Nanomaterials 2025, 15, 1528 13 of 21 Figure 9. SEM images (a,b) display specifics of the area near the notch of Cu/Nb-OL. The variationsof the surfaces around the notch (protrusion along RD) are attributed to the interfacial shear alongthe RD. Rotation of the nanolayers on the planes of the interfaces was evident, especially in (b). Thehorizontal direction here is TD, and the orthogonal direction into the image plane here is RD.Also, the load–displacement plot of the 63 nm Cu/Nb-OL TD microbeam is plotted inFigure 8. The load rises with the applied displacement of the indenter tip into the samplemainly due to the applied loading configuration (see Figure 6). This applied loadingconfiguration generates a combination of local bending moment and bending load whichconstrains the interfacial shear along the TD (constrains the flow of dislocation alongthe TD), while promoting the interfacial shear along the RD and hence local interfacialrotation (see Figure 9). Later, with an increase in the applied displacement, due to the localbending moment, the motion of dislocations is further constrained along the TD, resulting indislocation pileup and subsequent crack propagation. The crack propagation consequentlyresults in shear band formation and eventual failure of the beam. Such behavior wasalso observed during in situ TEM tensile tests of ARB 63 nm Cu/Nb nanolaminate byLiu et al. [50].The observations above were representative of a few microbeam bending experiments(3–5 microbeams) that we had conducted for each Cu/Nb-NOL and Cu/Nb-OL in thecurrent work. As discussed previously in Section 3.1, we were limited by the prohibitivelyhigh cost of the microbeam fabrication (using a focused ion beam (FIB)). However, withinour limited number of experiments, we observed the above phenomena (in Cu/Nb-NOLas well as in Cu/Nb-OL) very consistently (only differed in the extent)—both in thenotch widening and in the observation of interfacial rotation. We believe that the aboveexperiments provide an important piece of the interfacial plasticity mechanisms in Cu/NbARB nanolaminates and thus key insights which could lead to the full understanding of thedislocation processes and their interaction with the interfaces both along the TD and RDunder different loading configurations/conditions. The interfacial plasticity mechanism isan essential key for interface engineering and design for enhanced structural mechanicalperformance of nanomaterials in general and Cu/Nb nanolaminate-based mechanicalcomponents/structures specifically.4. DiscussionHere, we discussed the experimental observations of interfacial plasticity mechanismsof Cu/Nb ARB nanolaminate samples particularly along the RD which is otherwise theoret-ically restricted due to high interfacial shear strength documented for the similar materialsscheme [2,6,8,51–53], with our own work [15,16]. Typically, in Cu/Nb nanolaminates,owing to the lattice mismatch amongst copper and niobium, misfit dislocations are formedNanomaterials 2025, 15, 1528 14 of 21at the copper–niobium interface and reduce the overall strain energy. The interaction ofthese misfit dislocations with dislocations generated by externally applied loads governsthe localized interfacial plasticity of the Cu/Nb nanolaminate involving the plasticity inthe bulk nanolayers. When incoming dislocations due to externally applied load reach theinterface, they can interact with the misfit dislocations through the creation, adsorption,or annihilation of dislocations. The occurrence and extent of these dislocation-mediatedplasticity mechanisms depend on the specific characteristics of the misfit dislocations, suchas their Burgers vector, line direction, and the crystallographic orientation of the sample in-volved. As the ARB Cu/Nb nanolaminates have distinct crystallographic orientation alongthe TD and RD, they depict distinct behavior of dislocation-mediated plasticity mechanismsalong the TD and RD. The interfacial anisotropy of Cu/Nb nanolaminate has been exten-sively studied by means of experimental procedures [44,54–56] as well as computationalmodeling [8,14,41,57,58]. Molecular dynamic simulations by Demkowicz et al. [14] andWang et al. [41] showed that, for ARB Cu/Nb nanolaminate, the interface shear strength isevidently infinite along the RD whereas it has a finite value along the TD. From here on,we stress our discussion on the importance of crystallographic orientation along the RDand its effect on the extent of interfacial rotation as observed in the present study.It has been well documented in the literature from molecular dynamic simulationsperformed by Demkowicz et al. [14] and Wang et al. [41] that the interface shear strengthis theoretically infinite to start interfacial shear along the RD [41,59]. Therefore, in siturectangular micropillar compression experiments (with the Cu/Nb interfaces at ±45◦ fromthe loading axis) were performed along the RD and TD to directly probe the interfaceshear strength required to initiate dislocation-mediated plasticity along the RD and TD.These experiments serve as direct experimental verification of anisotropy in interfaceshear strength of the ARB Cu/Nb nanolaminate [14] explained in detail in Section 3.1.The in situ rectangular micropillar compression experiments clearly show that, althoughprohibited, interfacial shear could occur along the RD although accompanied by localplastic instabilities inducing other interface-mediated plasticity mechanisms, includinginterface rotation.In the instance of microbeam bending, both Cu/Nb-NOL and Cu/Nb-OL exhibitnotch widening along the TD, although the degree of the notch widening was higher inCu/Nb-NOL (~900%) compared to Cu/Nb-OL (with offset loading), which was ~400% ofthe original notch width. The notch widening could be the result of both interfacial plasticityas well as bulk plasticity (within the nanolayers). Nevertheless, due to the extensive plasticdeformation suffered by the constituent metal (especially Cu) in reduction from millimeterto nanometer scales during the ARB fabrication process, the available plasticity of thebulk material (within the nanolayers) is limited. Therefore, notch widening is primarilybecause of localized interfacial plasticity. This achievable interfacial shear strength alongthe TD (~1.2 GPa [14]) leads to interfacial shear in repsonse to the stress fields produced byan adjoining lattice dislocation under applied external load and induces the dislocationinside the interface [9,46]. Subsequently, the inward dislocation (regardless of the sign) getsabsorbed and then spreads its core in the interior of the interface plane, and it will contributeto interfacial plasticity and thereby notch widening. To go across the interfaces, thesedislocations would require re-nucleation to glide on the outward-bound slip system, whichcould happen only when extreme external stresses are imposed. Consequently, the ARB63 nm Cu/Nb TD microbeam provides stronger resistance to slip transmission [9,23]. Thehigher notch widening for Cu/Nb-NOL is attributed to unrestricted interfacial plasticityalong the TD compared to Cu/Nb-OL (with offset loading) where offset loading inducedthe local bending moment, restricting interfacial shearing along the TD.Nanomaterials 2025, 15, 1528 15 of 21Due to this restriction, we also observed interfacial rotation along the RD as hasbeen described in Section 3.2.2 through in situ microbeam bending of the Cu/Nb ARBmicrobeam with offset loading (Cu/Nb-OL) (see Figure 9). It has been noted that, duringthis in situ microbeam bending of nanolaminate (Cu/Nb-OL), shear vectors were present inboth the RD and TD, providing the possibility of layer rotational deformation in plane. Inthe case of Cu/Nb-OL, due to the constraint to the flow of dislocations along the TD becauseof offset loading, interfacial shear along the TD was restricted whereas interfacial shearalong the RD was activated, resulting in interfacial rotation. The existence of interfacialrotation (from TD towards RD) was verified by analyzing the front and back surfaces ofthe notch after in situ beam bending testing (see Figure 9). In Figure 9, we observe layershave protruded along the RD, exhibiting non-uniformity across the nanolayers, particularlynearby the notch area. This deformation is ascribed to the interfacial rotation along the RDdue to the constrained interfacial shear along the TD owing to the local bending momentintentionally induced in the Cu/Nb-OL (see Section 3.2.2).It is worth noting here that interfacial rotation, a critical phenomenon observed inthe present study, was also evident in our previous work. Specifically, we noted thisbehavior in our earlier publication, Anwarali et al. [47]. In that study, we briefly mentionedinterface rotation observation but did not elaborate it as it was not the focus of that study.The present study allows us to discuss further and highlight this important observation.In Anwarali et al. [47], we demonstrated interfacial rotation during in situ microbeambending experiments for unintended offset loading. Figure 10 pertaining to the work inAnwarali et al. [47] clearly illustrates the rotational deformation depicted via red arrows.Figure 10. SEM images showing the microbeam bending experiment with the observation of interfacerotation as earlier reported [47]: (a) the microbeam bending with offset loading setup and (b) the ob-servation of interfacial rotation due to the mixed-mode loading. Due to offset loading, the microbeamis subjected to the mixture of Mode-I and Mode-II loading, which results in the interfacial rotation asshown in the figure [47]. Reproduced with permission from Elsevier.By presenting Figure 10 alongside Figures 7 and 9 as have been shown earlier in thepresent manuscript, we provide additional visual proof of interface rotation occurringduring in situ microbeam bending. This provides additional supporting evidence to ourargument that this phenomenon is consistent and repeatable across different experimentswith certain intended/unintended offset loading (offset loading introduces shear stresses atthe crack tip, leading to a mixed-mode condition), validating the reliability of our findings.By emphasizing our previous work, we underline the continuity and progression of ourresearch in this area, showcasing the consistency of our observations over several yearsand across multiple in situ microbeam bending experiments. This continuity underscoresNanomaterials 2025, 15, 1528 16 of 21the depth of our work in understanding interfacial behavior, reinforcing the credibility andsignificance of our contributions to the field.Also, offset loading in a clamped beam bending geometry significantly affects themode of fracture and, consequently, the fracture toughness of a material [60]. This effect iscrucial in understanding and predicting the behavior of materials under mixed-mode load-ing conditions, which often occur in practical engineering applications. Fracture toughnessis a material’s capability to constrain crack spread under an applied load. It is quantified interms of stress intensity factors, KI for Mode-I (opening mode) and KII for Mode-II (slidingmode) fractures. In Mode-I, the crack surfaces are pulled apart perpendicularly, while inMode-II, the surfaces slide over each other in a parallel direction to the crack front. Thefracture toughness for each mode, KIC for Mode-I and KIIC for Mode-II, depends on thestress state nearby the crack tip, which can be changed by changing the loading conditions.In situ clamped microbeam bending is being used for evaluating fracture toughness,particularly because it provides a stable configuration for crack growth/propagation. In thisgeometry, the microbeam is clamped at both ends, and a load is applied at a point betweenthe clamps (see Figure 11a). The position of the load relative to the notch (offset distance,Lo) determines the mixity of the loading mode—whether it induces Mode-I, Mode-II, orboth (mixed mode) (see Figure 11b). When the load is applied directly above the crack (i.e.,Lo = 0), the loading condition is symmetric with respect to the crack line. This configurationinduces pure Mode-I loading, where the crack faces are pulled apart perpendicularly to thecrack front. The stress intensity factor KI dominates, and the material’s fracture toughnessis governed primarily by KIc, the Mode-I fracture toughness.KI =PLBW2√πaYI(aW,L0L)(1)KI I =PLBW2√πaYI I(aW,L0L)(2)where P, W, B, L, L0 and a are external applied load, width of the microbeam, thickness of themicrobeam, length of the microbeam, offset distance, and initial crack length, respectively.YI and YI I are mode I and mode II normalized geometric stress intensity factors, respectively,a function of relative crack length (a/W) and relative offset (L0/L) [60].Figure 11. Schematic illustrations of the in situ clamped microbeam bending setup with offset loading:(a) in situ clamped microbeam bending having a centered notch with offset loading configuration,and (b) schematic illustrations of Mode-I (opening) and Mode-II (in-plane shear).Nanomaterials 2025, 15, 1528 17 of 21As the load application point is moved away from the notch line (i.e., Lo > 0), anasymmetry is introduced in the loading condition. The offset loading introduces shearstresses at the crack tip, leading to a mixed-mode condition. This condition combines Mode-I and Mode-II loading together, with the proportion of Mode-II increasing as Lo increases.The stress intensity factors KI and KI I now both contribute to the crack propagation. Thefracture toughness is no longer governed solely by KIc but KI Ic, the Mode-II fracturetoughness, depending on the degree of mode mixity.Finite element method simulations [60] were employed to quantify the effect ofoffset loading on the stress intensity factors KI and KI I (Figures 11 and 2 (Chapter 4:Clamped Beam Bending for Mixed Mode Fracture Toughness Measurements) [60] andEquations (1) and (2)). With increasing Lo, the normalized KI decreases, indicating a reduc-tion in Mode-I contribution. Simultaneously, the normalized KI I increases, indicating anincreasing Mode-II contribution. However, after a certain point KI I begins to decrease again,signifying the complex interplay between the two modes. The fracture toughness measuredunder mixed-mode conditions differs from that measured under pure Mode-I or Mode-IIconditions. The toughness is influenced by the ratio KI I/KI , which varies with Lo. Formaterials that exhibit higher resistance to Mode-II (shear) than Mode-I (tensile) loading, thefracture toughness increases with increasing Lo. However, the exact relationship dependson the material’s intrinsic properties and the specific geometry of the test specimen.Similarly, in the present study, due to offset loading, the Cu/Nb ARB microbeamexperiences a mixture of Mode-I and Mode-II loading (Figure 12 in Reference [60]). For thecurrent fixed beam bending with offset loading configuration, Lo/L = 0.1, L/W = 8, anda/W = 0.3 replicate the scenario in Figure 2b (Chapter 4: Clamped Beam Bending for MixedMode Fracture Toughness Measurements) [60], which clearly shows that KI is reducingwhereas KI I is increasing, depicting a mixed-mode configuration. Due to the presence ofMode-I, the crack surfaces are pulled apart perpendicularly, resulting in interfacial slidingand subsequent notch widening, while due to the presence of Mode-II, the layers of Cu/NbARB nanolaminate slide over each other in a parallel direction to the crack front, resultingin the interfacial rotation as visible in Figures 7, 9 and 10.Comparing the two small-scale mechanical testing experiments of the Cu/Nb ARBnanolaminate samples, in situ rectangular micropillar compression offers more straight-forward experimental evidence of interfacial rotation in Cu/Nb ARB nanolaminates. Inmicropillar compression experiments, there was only the nominal shear stress (as thedriving force of deformation) in either the RD or TD. The loading allows the deformationin the Cu/Nb ARB interfaces either in the RD or TD, practically unconstrained by anyother macroscale (i.e., non-material) conditions. Despite nominally only shear stress in therolling direction (RD), our RD micropillar compression experiments exhibit clear evidenceof interfacial shearing also in the TD, hence the observation of the interfacial rotation(from RD towards TD) of the Cu/Nb nanolayers. The interfacial rotation was also evidentin the in situ microbeam bending with offset loading (Cu/Nb-OL), but the microbeambending mode here could lead to relatively more constraints (compared to a rectangularmicropillar) and thus much more complex stress states, especially near the notch tip, as hasbeen discussed in Section 3.1.2.Nevertheless, the experimental evidence in this present study has clearly shown inter-facial rotation of Cu/Nb nanolayers in the Cu/Nb ARB nanolaminate samples, despite norotational driving force from the loading mode (especially in the micropillar compression).Our aim in the present manuscript is to report this basic phenomenon. Our experimentalresults in the present study on localized interfacial plasticity along the TD and RD inCu/Nb nanolaminates have revealed the complex interplay between dislocation-mediatedlocalized plasticity, interfacial structure, and applied loading configurations/conditions.Nanomaterials 2025, 15, 1528 18 of 21Interfacial rotation is certainly an interesting phenomenon, but more importantly, withoutthe full understanding of interfacial rotation and interface-mediated plasticity mechanismsmore generally, we could not carry out interface engineering and design for enhanced struc-tural mechanical performance of nanomaterials in general and Cu/Nb nanolaminate-basedmechanical components/structures specifically. We believe the present study providessufficient justification for further, more in-depth, and quantitative studies of the interfacerotation mechanism in Cu/Nb ARB nanolaminate materials.5. ConclusionsIn situ microbeam bending as well as in situ micropillar compression experiments ofCu/Nb nanolaminates have revealed evidence regarding the dislocation-mediated local-ized plasticity along the RD and TD and the possible impact on the design/production ofCu/Nb-nanolaminate-based components/structures subjected to one or a combination ofcompression, torsional, and bending loading configurations. Interfacial shear and rotationwere observed along both the RD and TD during in situ rectangular pillar compression, de-spite theoretical predictions suggesting that shear should be inhibited along the RD. Offsetin situ microbeam bending along the TD introduced a local bending moment which re-stricted interfacial shear along the TD but generated interfacial shear/rotation along the RD.These outcomes of the study could be very significant for advanced structural applications,as they offer understandings of applied loading configurations/conditions and their effecton the interfacial shear/rotation along the TD and RD for anisotropic Cu/Nb nanolaminate.The authors think that information regarding interfacial sliding/rotation under diverseloading configurations could help researchers in scheming workable design diagrams for agiven material system, which could motivate the design/fabrication of novel strong/toughassemblies sustainable under complex loading configurations/conditions for evolvingfunctionalities, like stretchable bimetallic conductors for innovative wearable devices.Author Contributions: Conceptualization, I.R., M.J., M.E. and A.S.B.; Methodology, R.S. (Rahul Sahay),I.R., F.B., K.Y., M.J., D.S., M.E. and N.R.; Software, N.R. and A.S.B.; Validation, R.S. (Rahul Sahay),C.H. and T.S.; Formal analysis, R.S. (Rahul Sahay), I.R., C.H., F.B., K.Y., T.S., M.J., R.S. (Rachel Speaks),D.S., K.L., M.E., N.R. and A.S.B.; Investigation, R.S. (Rahul Sahay), I.R., P.A., C.H., F.B., K.Y., T.S., D.S.and N.R.; Resources, C.H., M.J., R.S. (Rachel Speaks), K.L. and M.E.; Data curation, K.Y. and T.S.;Writing—original draft, R.S. (Rahul Sahay) and A.S.B.; Writing—review & editing, R.S. (Rahul Sahay),F.B., M.J., R.S. (Rachel Speaks), D.S., K.L., M.E. and A.S.B.; Visualization, R.S. (Rahul Sahay), P.A. andA.S.B.; Supervision, M.J., D.S., M.E., N.R. and A.S.B.; Project administration, R.S. (Rachel Speaks), K.L.and N.R.; Funding acquisition, M.J., K.L., M.E., N.R. and A.S.B. All authors have read and agreed tothe published version of the manuscript.Funding: A.S.B., M.E. and M.J. thankfully recognize the funding support via e-ASIA Joint ResearchProgram (JRP) in the field of “Materials Informatics and Advanced Material Research by UtilizingComputers”—titled “Data-Driven Design of Mechanical Properties in Metallic Layered Structures”between Japan, Singapore, and Indonesia, 2021–2024. JSPS KAKENHI Grant No. 23H04464 and JSTSICORP Grant No. JPMJSC21E1. M.J. recognizes support from A*STAR, Singapore via the StructuralMetals and Alloys Programme (Grant No. A18B1b0061). A.S.B., R.S. and I.R. acknowledge co-fundingprovided by the Agence Nationale de la Recherche (ANR) of the French government through theGrant ANR18-09CE-003801 (Street Art Nano—Enabling Stretchable Metallic Conductors throughAtomic Reconfigurations in FCC/BCC Nanolayers) and the National Research Foundation (NRF) ofthe Singaporean government through the Grant NRF2018-NRF-ANR042 (Street Art Nano—EnablingStretchable Metallic Conductors through Atomic Reconfigurations in FCC/BCC Nanolayers). A.S.B.,R.S. and D.S. thankfully recognize the funding from Oregon Renewable Energy Center (OREC) underResearch Grant No. OREC2023/060/WIND. A.S.B. and K.L. also thankfully recognize the fundingfrom Oregon Renewable Energy Center (OREC) under Research Grant No. OREC2024/060/HPPT.Nanomaterials 2025, 15, 1528 19 of 21Data Availability Statement: The original contributions presented in this study are included in thearticle. Further inquiries can be directed to the corresponding author(s).Acknowledgments: The authors would like to thank Nathan Mara, University of Minnesota, USA andIrene Beyerlein, University of California, Santa Barbara for providing the Cu/Nb ARB samples. Wewould also like to acknowledge the assistance and expertise of Liu Qing of Temasek Labs, NanyangTechnological University (NTU), Singapore with the production of the microbeams for the tests.Conflicts of Interest: The authors declare no conflict of interest.References1. Holleck, H.; Schier, V. Multilayer PVD coatings for wear protection. Surf. Coat. Technol. 1995, 76, 328–336. [CrossRef]2. Beyerlein, I.J.; Mara, N.A.; Carpenter, J.S.; Nizolek, T.; Mook, W.M.; Wynn, T.A.; McCabe, R.J.; Mayeur, J.R.; Kang, K.; Zheng, S.;et al. Interface-driven microstructure development and ultra high strength of bulk nanostructured Cu-Nb multilayers fabricatedby severe plastic deformation. J. Mater. Res. 2013, 28, 1799–1812. [CrossRef]3. Hoagland, R.G.; Mitchell, T.E.; Hirth, J.P.; Kung, H. On the strengthening effects of interfaces in multilayer fee metallic composites.Philos. Mag. A 2002, 82, 643–664. [CrossRef]4. Han, W.; Cerreta, E.; Mara, N.; Beyerlein, I.; Carpenter, J.; Zheng, S.; Trujillo, C.; Dickerson, P.; Misra, A. Deformation and failureof shocked bulk Cu–Nb nanolaminates. Acta Mater. 2014, 63, 150–161. [CrossRef]5. Misra, A.; Demkowicz, M.J.; Zhang, X.; Hoagland, R.G. The radiation damage tolerance of ultra-high strength nanolayeredcomposites. JOM 2007, 59, 62–65. [CrossRef]6. Demkowicz, M.; Wang, Y.; Hoagland, R.; Anderoglu, O. Mechanisms of He escape during implantation in CuNb multilayercomposites. Nucl. Instrum. Methods Phys. Res. Sect. B Beam Interact. Mater. Atoms 2007, 261, 524–528. [CrossRef]7. Lim, S.; Rollett, A. Length scale effects on recrystallization and texture evolution in Cu layers of a roll-bonded Cu–Nb composite.Mater. Sci. Eng. A 2009, 520, 189–196. [CrossRef]8. Beyerlein, I.; Demkowicz, M.; Misra, A.; Uberuaga, B. Defect-interface interactions. Prog. Mater. Sci. 2015, 74, 125–210. [CrossRef]9. Wang, J.; Misra, A. An overview of interface-dominated deformation mechanisms in metallic multilayers. Curr. Opin. Solid StateMater. Sci. 2011, 15, 20–28. [CrossRef]10. Shao, S.; Medyanik, S.N. Dislocation–interface interaction in nanoscale fcc metallic bilayers. Mech. Res. Commun. 2010, 37, 315–319.[CrossRef]11. Zhang, R.; Germann, T.; Wang, J.; Liu, X.-Y.; Beyerlein, I. Role of interface structure on the plastic response of Cu/Nb nanolaminatesunder shock compression: Non-equilibrium molecular dynamics simulations. Scr. Mater. 2013, 68, 114–117. [CrossRef]12. Radchenko, I.; Zhu, W.; Qing, L.; Navarro, E.; Sahay, R.; Lee, P.S.; Chen, K. Interface Rotation in Cu/Nb Accumulative RollBonded (ARB) Nanolaminates. SSRN Electron. J. 2022. [CrossRef]13. Sahay, R.; Budiman, A.S.; Aziz, I.; Navarro, E.; Escoubas, S.; Cornelius, T.W.; Gunawan, F.E.; Harito, C.; Lee, P.S.; Thomas, O.; et al.Crystallographic Anisotropy Dependence of Interfacial Sliding Phenomenon in a Cu(16)/Nb(16) ARB (Accumulated RollingBonding) Nanolaminate. Nanomaterials 2022, 12, 308. [CrossRef]14. Demkowicz, M.; Thilly, L. Structure, shear resistance and interaction with point defects of interfaces in Cu–Nb nanocompositessynthesized by severe plastic deformation. Acta Mater. 2011, 59, 7744–7756. [CrossRef]15. Ali, H.P.A.; Budiman, A. Advances in In situ microfracture experimentation techniques: A case of nanoscale metal–metalmultilayered materials. J. Mater. Res. 2019, 34, 1449–1468. [CrossRef]16. Radchenko, I.; Anwarali, H.P.; Tippabhotla, S.K.; Budiman, A.S. Effects of interface shear strength during failure of sem-icoherentmetal–metal nanolaminates: An example of accumulative roll-bonded Cu/Nb. Acta Mater. 2018, 156, 125–135. [CrossRef]17. Ali, H.P.A.; Radchenko, I.; Li, N.; Budiman, A. The roles of interfaces and other microstructural features in Cu/Nb nanolayersas revealed by in situ beam bending experiments inside an scanning electron microscope (SEM). Mater. Sci. Eng. A 2018, 738,253–263. [CrossRef]18. Ali, H.P.A.; Radchenko, I.; Li, N.; Budiman, A. Effect of multilayer interface through in situ fracture of Cu/Nb and Al/Nb metallicmultilayers. J. Mater. Res. 2019, 34, 1564–1573. [CrossRef]19. Zheng, S.; Wang, J.; Carpenter, J.; Mook, W.; Dickerson, P.; Mara, N.; Beyerlein, I. Plastic instability mechanisms in bimetallicnanolayered composites. Acta Mater. 2014, 79, 282–291. [CrossRef]20. Saito, Y.; Utsunomiya, H.; Tsuji, N.; Sakai, T. Novel ultra-high straining process for bulk materials—Development of theaccumulative roll-bonding (ARB) process. Acta Mater. 1999, 47, 579–583. [CrossRef]21. Wang, J.; Hoagland, R.; Liu, X.; Misra, A. The influence of interface shear strength on the glide dislocation–interface interactions.Acta Mater. 2011, 59, 3164–3173. [CrossRef]22. Wang, J.; Zhou, Q.; Shao, S.; Misra, A. Strength and plasticity of nanolaminated materials. Mater. Res. Lett. 2017, 5, 1–19.[CrossRef]https://doi.org/10.1016/0257-8972(95)02555-3https://doi.org/10.1557/jmr.2013.21https://doi.org/10.1080/01418610208243194https://doi.org/10.1016/j.actamat.2013.10.019https://doi.org/10.1007/s11837-007-0120-6https://doi.org/10.1016/j.nimb.2007.04.110https://doi.org/10.1016/j.msea.2009.05.020https://doi.org/10.1016/j.pmatsci.2015.02.001https://doi.org/10.1016/j.cossms.2010.09.002https://doi.org/10.1016/j.mechrescom.2010.03.007https://doi.org/10.1016/j.scriptamat.2012.09.022https://doi.org/10.2139/ssrn.4101121https://doi.org/10.3390/nano12030308https://doi.org/10.1016/j.actamat.2011.09.004https://doi.org/10.1557/jmr.2019.75https://doi.org/10.1016/j.actamat.2018.06.023https://doi.org/10.1016/j.msea.2018.09.094https://doi.org/10.1557/jmr.2018.449https://doi.org/10.1016/j.actamat.2014.07.017https://doi.org/10.1016/S1359-6454(98)00365-6https://doi.org/10.1016/j.actamat.2011.01.056https://doi.org/10.1080/21663831.2016.1225321Nanomaterials 2025, 15, 1528 20 of 2123. Wang, J.; Misra, A.; Hoagland, R.; Hirth, J. Slip transmission across fcc/bcc interfaces with varying interface shear strengths. ActaMater. 2012, 60, 1503–1513. [CrossRef]24. Misra, A.; Hirth, J.P.; Kung, H. Single-dislocation-based strengthening mechanisms in nanoscale metallic multilayers. Philos. Mag.A 2002, 82, 2935–2951. [CrossRef]25. Zheng, S.; Beyerlein, I.J.; Carpenter, J.S.; Kang, K.; Wang, J.; Han, W.; Mara, N.A. High-strength and thermally stable bulknanolayered composites due to twin-induced interfaces. Nat. Commun. 2013, 4, 1696. [CrossRef] [PubMed]26. Nix, W.D.; Gao, H. Indentation size effects in crystalline materials: A law for strain gradient plasticity. J. Mech. Phys. Solids 1998,46, 411–425. [CrossRef]27. Wu, K.; Zhang, J.; Zhang, P.; Wang, Y.; Liu, G.; Zhang, G.; Sun, J. Fracture behavior and adhesion energy of nanostructuredCu/Mo multilayer films. Mater. Sci. Eng. A 2014, 613, 130–135. [CrossRef]28. Greer, J.R.; Oliver, W.C.; Nix, W.D. Size dependence of mechanical properties of gold at the micron scale in the absence of straingradients. Acta Mater. 2005, 53, 1821–1830. [CrossRef]29. Greer, J.R.; De Hosson, J.T.M. Plasticity in small-sized metallic systems: Intrinsic versus extrinsic size effect. Prog. Mater. Sci. 2011,56, 654–724. [CrossRef]30. Mara, N.A.; Bhattacharyya, D.; Dickerson, P.; Hoagland, R.G.; Misra, A. Deformability of ultrahigh strength 5 nm Cu/Nbnanolayered composites. Appl. Phys. Lett. 2008, 92, 231901. [CrossRef]31. Zhang, H.; Schuster, B.; Wei, Q.; Ramesh, K. The design of accurate micro-compression experiments. Scr. Mater. 2006, 54, 181–186.[CrossRef]32. Li, N.; Mara, N.; Wang, J.; Dickerson, P.; Huang, J.; Misra, A. Ex situ and in situ measurements of the shear strength of interfacesin metallic multilayers. Scr. Mater. 2012, 67, 479–482. [CrossRef]33. Jaya, B.N.; Kirchlechner, C.; Dehm, G. Can microscale fracture tests provide reliable fracture toughness values? A case study insilicon. J. Mater. Res. 2015, 30, 686–698. [CrossRef]34. Mayer, C.; Yang, L.; Singh, S.; Llorca, J.; Materialia, J.L.-A.; Shen, Y.L.; Chawla, N. Anisotropy, size, and aspect ratio effects onmicropillar compression of AlSiC nanolaminate composites. Acta Mater. 2016, 114, 25–32. [CrossRef]35. Zeng, Z.; Xiao, Y.; Wheeler, J.M.; Tan, J.-C. In situ micropillar compression of an anisotropic metal-organic framework singlecrystal. Commun. Chem. 2023, 6, 63. [CrossRef] [PubMed]36. Maeder, X.; Mook, W.; Niederberger, C.; Michler, J. Quantitative stress/strain mapping during micropillar compression. Philos.Mag. 2011, 91, 1097–1107. [CrossRef]37. Dimiduk, D.M.; Woodward, C.; LeSar, R.; Uchic, M.D. Scale-free intermittent flow in crystal plasticity. Science 2006, 312, 1188–1190.[CrossRef] [PubMed]38. Weygand, D.; Poignant, M.; Gumbsch, P.; Kraft, O. Three-dimensional dislocation dynamics simulation of the influence of samplesize on the stress–strain behavior of fcc single-crystalline pillars. Mater. Sci. Eng. A 2008, 483–484, 188–190. [CrossRef]39. Zhou, J.; Averback, R.; Bellon, P. Stability and amorphization of Cu–Nb interfaces during severe plastic deformation: Moleculardynamics simulations of simple shear. Acta Mater. 2014, 73, 116–127. [CrossRef]40. Zheng, S.; Beyerlein, I.; Wang, J.; Carpenter, J.; Han, W.; Mara, N. Deformation twinning mechanisms from bimetal interfaces asrevealed by in situ straining in the TEM. Acta Mater. 2012, 60, 5858–5866. [CrossRef]41. Wang, J.; Kang, K.; Zhang, R.F.; Zheng, S.J.; Beyerlein, I.J.; Mara, N.A. Structure and property of interfaces in ARB Cu/Nblaminated composites. JOM 2012, 64, 1208–1217. [CrossRef]42. Anderson, P.M.; Bingert, J.F.; Misra, A.; Hirth, J.P. Rolling textures in nanoscale Cu/Nb multilayers. Acta Mater. 2003, 51,6059–6075. [CrossRef]43. Anderson, P.; Foecke, T.; Hazzledine, P. Dislocation-based deformation mechanisms in metallic nanolaminates. MRS Bull. 1999,24, 27–33. [CrossRef]44. Nizolek, T.; Beyerlein, I.J.; Mara, N.A.; Avallone, J.T.; Pollock, T.M. Tensile behavior and flow stress anisotropy of accumulativeroll bonded Cu-Nb nanolaminates. Appl. Phys. Lett. 2016, 108, 051903. [CrossRef]45. Wang, J.; Hoagland, R.; Hirth, J.; Misra, A. Atomistic simulations of the shear strength and sliding mechanisms of copper–niobiuminterfaces. Acta Mater. 2008, 56, 3109–3119. [CrossRef]46. Wang, J.; Hoagland, R.G.; Hirth, J.P.; Misra, A. Atomistic modeling of the interaction of glide dislocations with “weak” interfaces.Acta Mater. 2008, 56, 5685–5693. [CrossRef]47. Budiman, A.S.; Sahay, R.; Ali, H.P.A.; Tippabhotla, S.K.; Radchenko, I.; Raghavan, N. Interface-mediated plasticity and fracture innanoscale Cu/Nb multilayers as revealed by in situ clamped microbeam bending. Mater. Sci. Eng. A 2021, 803, 140705. [CrossRef]48. Pilkey, W.D.; Pilkey, D.F. Peterson’s Stress Concentration Factors; John Wiley & Sons: Hoboken, NJ, USA, 2008.49. Li, Y.; Zhou, Q.; Zhang, S.; Huang, P.; Xu, K.; Wang, F.; Lu, T. On the role of weak interface in crack blunting process in nanoscalelayered composites. Appl. Surf. Sci. 2018, 433, 957–962. [CrossRef]50. Liu, Z.; Monclús, M.; Yang, L.; Castillo-Rodríguez, M.; Molina-Aldareguía, J.; Llorca, J. Tensile deformation and fracturemechanisms of Cu/Nb nanolaminates studied by in situ TEM mechanical tests. Extrem. Mech. Lett. 2018, 25, 60–65. [CrossRef]https://doi.org/10.1016/j.actamat.2011.11.047https://doi.org/10.1080/01418610208239626https://doi.org/10.1038/ncomms2651https://www.ncbi.nlm.nih.gov/pubmed/23591863https://doi.org/10.1016/S0022-5096(97)00086-0https://doi.org/10.1016/j.msea.2014.06.083https://doi.org/10.1016/j.actamat.2004.12.031https://doi.org/10.1016/j.pmatsci.2011.01.005https://doi.org/10.1063/1.2938921https://doi.org/10.1016/j.scriptamat.2005.06.043https://doi.org/10.1016/j.scriptamat.2012.06.008https://doi.org/10.1557/jmr.2015.2https://doi.org/10.1016/j.actamat.2016.05.018https://doi.org/10.1038/s42004-023-00858-whttps://www.ncbi.nlm.nih.gov/pubmed/37016101https://doi.org/10.1080/14786435.2010.505178https://doi.org/10.1126/science.1123889https://www.ncbi.nlm.nih.gov/pubmed/16728635https://doi.org/10.1016/j.msea.2006.09.183https://doi.org/10.1016/j.actamat.2014.03.055https://doi.org/10.1016/j.actamat.2012.07.027https://doi.org/10.1007/s11837-012-0429-7https://doi.org/10.1016/S1359-6454(03)00428-2https://doi.org/10.1557/S0883769400051514https://doi.org/10.1063/1.4941043https://doi.org/10.1016/j.actamat.2008.03.003https://doi.org/10.1016/j.actamat.2008.07.041https://doi.org/10.1016/j.msea.2020.140705https://doi.org/10.1016/j.apsusc.2017.10.002https://doi.org/10.1016/j.eml.2018.10.007Nanomaterials 2025, 15, 1528 21 of 2151. Carpenter, J.S.; McCabe, R.J.; Zheng, S.J.; Wynn, T.A.; Mara, N.A.; Beyerlein, I.J. Processing parameter influence on texture andmicrostructural evolution in Cu-Nb multilayer composites fabricated via accumulative roll bonding. Metall. Mater. Trans. A 2014,45, 2192–2208. [CrossRef]52. Hattar, K.; Misra, A.; Dosanjh, M.R.F.; Dickerson, P.; Robertson, I.M.; Hoagland, R.G. Direct observation of crack propagation incopper–niobium multilayers. J. Eng. Mater. Technol. 2012, 134, 021014. [CrossRef]53. Beyerlein, I.; Mayeur, J.; McCabe, R.; Zheng, S.; Carpenter, J.; Mara, N. Influence of slip and twinning on the crystallographicstability of bimetal interfaces in nanocomposites under deformation. Acta Mater. 2014, 72, 137–147. [CrossRef]54. Demkowicz, M.J.; Beyerlein, I.J. The effects of nanoscale confinement on the behavior of metal laminates. Scr. Mater. 2020, 187,130–136. [CrossRef]55. Chen, T.; Yuan, R.; Beyerlein, I.J.; Zhou, C. Predicting the size scaling in strength of nanolayered materials by a discrete slip crystalplasticity model. Int. J. Plast. 2020, 124, 247–260. [CrossRef]56. Dong, S.; Chen, T.; Huang, S.; Li, N.; Zhou, C. Thickness-dependent shear localization in Cu/Nb metallic nanolayered composites.Scr. Mater. 2020, 187, 323–328. [CrossRef]57. Budiman, A.S.; Han, S.-M.; Li, N.; Wei, Q.-M.; Dickerson, P.; Tamura, N.; Kunz, M.; Misra, A. Plasticity in the nanoscale Cu/Nbsingle-crystal multilayers as revealed by synchrotron Laue x-ray microdiffraction. J. Mater. Res. 2012, 27, 599–611. [CrossRef]58. Misra, A.; Demkowicz, M.J.; Wang, J.; Hoagland, R.G. The multiscale modeling of plastic deformation in metallic nanolayeredcomposites. JOM 2008, 60, 39–42. [CrossRef]59. Mara, N.A.; Bhattacharyya, D.; Hirth, J.P.; Dickerson, P.; Misra, A. Mechanism for shear banding in nanolayered composites. Appl.Phys. Lett. 2010, 97, 021909. [CrossRef]60. Jonnalagadda, K.; Alankar, A.; Balila, N.J.; Bhandakkar, T. Advances in Structural Integrity: Structural Integrity over Multiple LengthScales; Springer: Berlin/Heidelberg, Germany, 2022.Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individualauthor(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury topeople or property resulting from any ideas, methods, instructions or products referred to in the content.https://doi.org/10.1007/s11661-013-2162-4https://doi.org/10.1115/1.4005953https://doi.org/10.1016/j.actamat.2014.03.041https://doi.org/10.1016/j.scriptamat.2020.05.057https://doi.org/10.1016/j.ijplas.2019.08.016https://doi.org/10.1016/j.scriptamat.2020.06.049https://doi.org/10.1557/jmr.2011.421https://doi.org/10.1007/s11837-008-0047-6https://doi.org/10.1063/1.3458000 Introduction  Materials and Methods  Assembly of ARB Cu/Nb Nanolaminates  Fabrication of Cu/Nb ARB Rectangular Micropillar Samples  Fabrication of Cu/Nb ARB Microbeam Samples  In Situ Microbeam Bending  Experimental Results  In Situ Rectangular Micropillar Compression  Interface Shear in the TD Rectangular Micropillar Compression  Interface Plasticity in the RD Rectangular Micropillar Compression  In Situ Cu/Nb ARB Microbeam Bending  In Situ Microbeam Bending of Cu/Nb ARB Microbeam with No Offset Loading (Cu/Nb-NOL)  In Situ Microbeam Bending Experiments of Cu/Nb ARB Microbeam with Offset Loading (Cu/Nb-OL)  Discussion  Conclusions  References