Tailoring Fast Directional Mass Transport of Nano-Confined AgâCu

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Surfaces, Interfaces, and Applications

Tailoring Fast Directional Mass Transport of Nano-Confined AgCu Alloys upon Heating: Effect of the AlN Barrier Thickness Vicente Araullo-Peters, Claudia Cancellieri, Mirco Chiodi, Jolanta Janczak-Rusch, and Lars P.H. Jeurgens ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b19091 • Publication Date (Web): 17 Jan 2019 Downloaded from http://pubs.acs.org on January 21, 2019

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ACS Applied Materials & Interfaces

Tailoring Fast Directional Mass Transport of NanoConfined Ag-Cu Alloys upon Heating: Effect of the AlN Barrier Thickness Vicente Araullo-Peters,† Claudia Cancellieri, Mirco Chiodi, Jolanta JanczakRusch, Lars P.H. Jeurgens* Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Joining Technologies & Corrosion, Ueberlandstrasse 129, 8600 Duebendorf, Switzerland

ABSTRACT This study addresses the phase stability and atomic mobility of Ag-Cu40at% nano-alloys confined by AlN in a nanomultilayered configuration during thermal treatment. To this end, nanomultilayers (NMLs) with a fixed Ag-Cu40at% nanolayer thickness of 8 nm and a variable AlN barrier nanolayer thickness of 4 nm, 8 nm or 10 nm were deposited by magnetron sputtering on sapphire substrates and subsequently isothermally annealed for 5 or 20 minutes in air up in the range of 200 - 500 ˚C. The microstructure of the asdeposited and heat-treated NMLs were analysed by XRD, SEM, TEM and EDS. Annealing of the thicker AlN barrier layers at T > 300 °C leads to the formation of an interconnected network of line-shaped Cu(O) protrusions on the annealed NML surface. The well-defined outflow pattern of Cu(O) originates from the thermally-induced surface cracking of the top AlN barriers with subsequent fast mass transport of Cu along the AgCu/AlN interfaces towards the surface cracks. The thinnest (i.e. 4 nm thick) AlN barrier layers exhibit a relatively open grain boundary structure and act as nanoporous membranes upon heating, resulting in the formation of a dense and homogenous -1ACS Paragon Plus Environment

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distribution of Cu(O) and Ag droplets on the NML surface. These findings demonstrate that the microstructure (i.e. layer thicknesses, interface coherency and texture) of hybrid nanolaminates can be tuned to provide defined pathways for fast, directional transport of the confined metal to the surface at relatively low temperatures, which might open new routes for low-temperature bonding of micro- and nano-scaled systems. KEYWORDS: Multilayer thin films, nanoconfinement, interface migration, thermally activated processes, nano-joining

1.

INTRODUCTION

The joining of dissimilar materials (e.g. metals, semiconductors, ceramics, glass, polymers, plastics) remains a non-trivial and ever-present challenge for the design and engineering of multi-material assemblies.1-2 The continuous miniaturization and diversification of functional devices and sensing components calls for novel hybrid joining technologies, which enable bonding of different materials as well as system assembly and packaging at ever-lower temperatures with ever-higher alignment accuracy.3 At the same time, the use of lead-containing solders and the addition of chemical solvents (e.g. chemical flux) for soldering should be avoided.3 Moreover, the joints and interconnections should be tuned not only for their (thermo)mechanical strength, but in particular also for their thermal and electrical conductivity (e.g. improving the contact resistance and/or enabling more-efficient cooling of miniaturized high-power-density devices).3 Silver and copper are the most commonly used metals in conventional joining technologies and both exhibit excellent thermal and electrical conductivities. Eutectic Ag-40 at.% Cu fillers have a melting temperature of 779 ˚C, well below the bulk melting temperatures of pure silver (i.e. 962˚C) and pure copper (i.e. 1085˚C). Notably, Ag and Cu are practically immiscible below ~300 °C and, consequently, a large driving force for -2ACS Paragon Plus Environment

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phase segregation of (super)saturated Ag-Cu solid solutions exists below the eutectic melting temperature.4-5 The melting point of Ag-Cu alloys can be further lowered by the addition of a melting point depressant (e.g. indium, zinc, tin, phosphor, boron), which often deteriorates the joint performance (e.g. mechanical strength, corrosion resistance).1 In recent years, various novel bonding technologies have been envisaged as a replacement of conventional Cu- and Ag-based joining technologies.6-10 For example, novel solder pastes constituted of Ag, AgO or CuO nanoparticles (NP) (typically coated with an organic solvent) have been developed, which exhibit enhanced sintering kinetics at relatively low temperatures, as driven by very fast surface diffusion in possible combination with surface melting.6-7,11 Joining processes with Ag- and CuO-nanopastes can be performed at temperatures as low as 250 ˚C and 400 ˚C, respectively.6-7 However, broad application of nanopaste technologies suffers from handling, safety and processing issues, such as nanoparticle agglomeration, oxidation (ageing), sintering, organic residues and the need to apply large external pressures to reduce porosity in the joint zone. Surface-activated bonding (SAB) presents an alternative solution for lowtemperature bonding8-9,12-13 and also (i.e. as for nanopaste) exploits the intrinsically high reactivity of atomically-clean metal surfaces. For example, Cu–Cu direct bonding was achieved at 250 ˚C under high vacuum conditions with an external pressure of about 1 MPa by exploiting fast surface diffusion on Cu(111) planes in combination with surfacediffusion controlled creep.14 Analogously, pressure-less Ag-Ag direct bonding at 250 ˚C in air was achieved by maximizing the compressive growth stresses in the as-deposited Ag layers to drive abnormal grain growth at the contacting interface.8-9 Evidently, SAB bonding technologies are very sensitive to the cleanness and roughness of the bonding surfaces and thus require atomically controlled surface preparation methods, as well as high vacuum equipment for the bonding of oxidation-sensitive metals. -3ACS Paragon Plus Environment


Our research team envisages another class of nanostructured filler materials, namely hybrid nanolaminates, for selective bonding at reduced temperatures, which could tackle some of the aforementioned limitations of emerging nanojoining methods.2,5,15-18. The nanolaminates (also referred to as nanomultilayers; NMLs) are composed of a stack of alternating nanolayers (NLs) of a metal or an alloy (here: eutectic Ag-Cu40at.%) and a chemically inert barrier material (a nitride, oxide or refractory metal; here AlN). Such hybrid NML systems can be easily produced as a coating or as a foil by (reactive) magnetron sputtering with accurate control of the NL thickness, modulation periodicity, chemical composition, in-plane texture and growth stress5,15,18-19 Figs. 1a,b show crosssectional and planar SEM micrographs of the type of hybrid nanolaminate addressed in the present study, which is composed of alternating nanolayers of an Ag–Cu40at.% alloy and an AlN barrier.

Figure 1. Secondary electron image of (a) a cross section and (b) the surface of an as-deposited AgCu40at.%/AlN NML, as produced by magnetron sputtering on a sapphire substrate. The NML stack is constituted of a 10 nm bottom AlN buffer layer with 10 repetitions of a [AgCu10nm/AlN10nm] building block on top (dark grey: AlN, light grey: Ag-Cu).

The ultimate goal is to trigger the very fast, directional mass transport of the confined metal filler to the NML surface at relatively low temperatures, which may be exploited for localized chemical bonding of micro- and nano-scaled components.20-22 To this end, the inert barrier NLs (here: AlN, see Fig. 1) should confine the metal NLs only up to the point (i.e. temperature) where the diffusion of the confined metal atoms becomes triggered -4ACS Paragon Plus Environment

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(i.e. thermally activated). Hence, the barrier layers should be pre-designed to provide pathways (e.g. cracks, grain boundaries, structural defects or pores) for directed (or even patterned) outflow of the confined metal to the bonding surface for selective bonding. The present study reports important criteria for microstructural design of NML systems to achieve fast directional outflow of the confined metal to the NML surface upon heating in air. In particular, it is demonstrated that the thickness and morphology of the AlN barrier layers can be tuned such to provoke different modes of fast and directional mass transport of the confined metal to the NML surface at relatively low temperatures. Notably, the obtained fundamental knowledge on the phase stability and atomic mobility of confined solids in hybrid nanolaminates is not only of interest to the microand nano-joining community, but also highly relevant for many other application areas, such as hard coatings, optical filters, X-ray mirrors, reactive joining, energy storage, micro-electronics and plasmonics.23-25

2.

EXPERIMENTAL SECTION

Ag-Cu40at.%/AlN NML coatings were deposited on 2” epi-polished -Al2O3(0001) singlecrystalline wafer substrates by magnetron sputtering in a high vacuum chamber (base pressure < 110-8 mbar) from two confocally arranged unbalanced magnetrons equipped with 2" targets of pure Al (99.99% purity, as supplied by Kurt J. Lesker, USA) and of a bulk eutectic Ag-Cu40at.% alloy. Before insertion in the deposition apparatus, the sapphire substrates were ultrasonically cleaned using Acetone and Ethanol. Prior to deposition, possible surface contamination on the sapphire substrate (mostly adventitious C) was removed by Ar+ sputter cleaning for 5 min applying a RF Bias of 100 V. Further experimental details on the NML deposition process are given in Refs. 5,15. First, a 10 nm AlN barrier layer was deposited on the sputter-cleaned substrate. Next 20 repetitions of alternating Ag-Cu40at.% and AlN nanolayers (NLs) were deposited on top. In -5ACS Paragon Plus Environment


the present study, the thickness of the Ag-Cu40at.% NL was fixed at 8 nm, resulting in a cumulated Ag-Cu40at.% NL thickness of 20×8 = 160 nm. Only the AlN barrier layer thickness was varied for each NML deposition, selecting a constant thickness of either 4, 8 or 10 nm (resulting in cumulated barrier thickness of 10 nm + either 80, 160 or 200 nm, respectively). The NML specimens with barrier thicknesses of 4 nm, 8 nm and 10 nm and a fixed alloy layer thickness of 8 nm are further referred to as AgCu8nm|AlN4nm, AgCu8nm|AlN8nm and AgCu8nm|AlN10nm, respectively. As experienced from various trial depositions in combination with AFM surface scans (see Fig. S1 in the Supplementary Material) and cross-sectional TEM, a well-defined nanolaminated structure with an effective nano-confinement of the alloy is obtained for barrier thicknesses > 3 nm (see also).5,15 Therefore, only barrier layer thicknesses > 3 nm are considered in the study. Next, the thus-produced NML specimens were isothermally annealed in air for different annealing times in the range of 5 to 20 minutes and different annealing temperatures in the range of 200 - 500˚C. Near-isothermal annealing conditions were achieved by placing the specimen into a partial open cavity of a massive Cu block, which was positioned and preheated at the desired temperature in a conventional furnace (see Fig. S2 in the supplementary material). After each thermal treatment, the specimen was removed from the Cu block and allowed to cool in air. High-resolution SEM (HR-SEM) analysis of the as-deposited and annealed specimens was performed using a Hitachi S-4800 instrument equipped with a Bruker XFlash 6|60 energy dispersive X-ray detector. Surface FIB cross-sectional cuts were conducted using a FEI Helios 660 Dualbeam. A Jeol 2200 TEM and the Fei Helios 660 Dualbeam were used for STEM imaging with energy dispersive spectroscopy (EDS). A Discover (D8) X-ray diffractometer from Bruker, operated in Bragg-Brentano geometry, was applied to measure 2-theta (2) scans of the as-deposited and annealed specimens. In addition,

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pole figures of the Ag{111 and AlN{103} family of planes were recorded to investigate the in-plane and out-of-plane texture of the alloy and barrier NLs.

3.

RESULTS AND DISCUSSION

Ag-Cu40at.%/AlN NML with a fixed alloy NL thickness of 8 nm and a variable barrier NL thicknesses of 4 nm, 8 nm and 10 nm (as based on previous findings5,15,18) were deposited by magnetron sputtering on sapphire substrates and subsequently isothermally annealed in air at various temperatures in the range of 200 - 500˚C for durations of 5 and 20 minutes (See experimental Section). Microstructure of the as-deposited NMLs: Cross-sectional and planar HR-SEM micrographs of an as-deposited AgCu8nm|AlN10nm NML on the -Al2O3(0001) substrate are shown in Figs. 1a and 1b, respectively. The NML sequence, as composed of a 10 nm thick bottom AlN buffer layer with 10 repetitions of a [AgCu8nm/AlN10nm] building block on top (dark grey: AlN, light grey: AgCu), is very nicely resolved (Fig. 1a). As evidenced by diffraction analysis,5 the as-deposited Ag-Cu40at.% NLs are composed of a fcc matrix of Ag nano-grains, which are supersaturated by Cu (further denoted as Ag[Cu]), and possibly few smaller embedded Cu nano-grains. Hence the alloy NLs are metastable; segregation of Cu out of the supersaturated Ag[Cu] nano-grains phase (i.e. phase separation) becomes thermally activated at T  250 °C.5 The individual thicknesses of the AgCu8nm and AlN10nm NLs are nearly constant, which in turn results in a very smooth and uniform (i.e. featureless) NML surface morphology: compare Figs. 1a and 1b, respectively. This suggests a layer-by-layer-like growth mode for the co-deposition of Ag and Cu on AlN using the here-applied Ag[Cu] and AlN deposition conditions (see Experimental Section). On the contrary, deposition of pure Ag on AlN under similar conditions (i.e. in the absence of Cu) results in a progressive roughening of the successively deposited NLs towards the NML surface (as attributed to the poor wetting of Ag on AlN).15 -7ACS Paragon Plus Environment


Fig. 2 shows the XRD analysis of the as-deposited AgCu8nm|AlN10nm NML. The pole figures of the Ag(111) and AlN(10-13) family of planes confirm that the alloy and barrier layers are nano-crystalline and both possess a strong in-plane and out-of-plane texture. The out-of-plane texture originates from the preferred growth of the Ag[Cu] and the AlN nano-grains along the Ag[111] and AlN[0001] directions, respectively (corresponding to the energetically-preferred planes of Ag and AlN being parallel to the film surface).5,26 The strong in-plane texture of the six-fold symmetric Ag(111) and the hexagonallypacked AlN(0001) planes stems from the orientation relationships (OR) between the substrate, the AlN barrier and the Ag[Cu] nano-grains and results in a six-fold symmetry in the respective pole figures (see Figs. 2a and 2b). Recorded Phi-scans (cf. Fig. 2c) evidence an OR between the -Al2O3(0001) substrate, the AlN barrier and the Ag[Cu] NL according to Ag{111}AlN{0001}Al2O3{0001}. The established OR between AlN and Ag[Cu] is in accordance with previous findings5,15,26 and is inherited (albeit slightly weakened5,15,19) throughout the multilayer stack.

Figure 2. XRD pole figures for (a) the Ag(111) and (b) the AlN(0103) family of planes, as recorded from the AgCu8nm|AlN10nm NML in the as-deposited state. (c) Phi-scans of the α-Al203(2-1-13), Ag(111) and AlN(10-13) reflections.

Strikingly, the OR between Ag[Cu] and AlN is more pronounced than that between pure Ag and AlN for comparable NML systems deposited under similar conditions (as -8ACS Paragon Plus Environment

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evidenced from the much sharper and stronger peak intensities in the recorded phiscans of the AgCu40at%/AlN NML).15 Hence the dissolution of Cu into Ag during the codeposition not only promotes layer-by-layer growth (see above), but also enhances the coherency of the AgAlN interface (associated with a tensile strain contribution due to the lattice mismatch5). Notably, the in-plane coherency and associated in-plane texture of the Ag[Cu] NMLs (as evidenced from the measured pole figures and phi-scans) decreases with decreasing AlN barrier thickness. Topographic analysis of the NML surface (i.e. the top AlN barrier layer) by highresolution AFM (see Fig. S1 in the Supplementary Material) evidences that the thickest (i.e. 10 nm thick) AlN barrier layers are relatively compact and of uniform thickness, whereas the 4nm-thick AlN barrier layers appear less compact due to the presence of tiny nanopores (see green arrows in Fig. S1b), which are rather randomly distributed across the NML surface. These findings are characteristic for the initial stages of thin film growth by physical vapour deposition methods, as associated with the gradual coalescence of initially isolated islands (nucleated on the surface) into a laterally-closed polycrystalline thin film.27-29 The coalescence of individual AlN islands into a compact AlN barrier layer is not yet fully completed for barrier thicknesses as small as 4 nm, resulting in a relatively open grain boundary structure at locations where AlN nanograins have not yet fully coalesced (evidenced as nanopores in Fig. S1b).29 The lower degree of coalescence of the thinnest AlN barriers is accompanied by decrease of the in-plane coherency of the Ag[Cu]AlN interface, as evidenced by XRD (see above). Nevertheless, as shown by cross-sectional HRTEM analysis, the 4-nm-thick AlN barrier layers still have a relatively uniform thickness and "appear" laterally continuous in the 2-dimensional (2D) TEM projection of the 3-dimensional (3D) AlN nanolayers.1 Hence, the thinnest 4-

In this regard, it is emphasized that an open grain-boundary structure due to incomplete coalescence of AlN nanograins cannot be visualized by cross-sectional TEM analysis, since the HRTEM image is a 2D 1

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nm-thick AlN barrier layers are still capable of imposing an effective nano-confinement of the alloy; however, their thermomechanical response upon heating strikingly differs from the thicker AlN barrier layers, as discussed in the following. Microstructure of the annealed NMLs: Planar-view HR-SEM analysis of the AgCu8nm|AlN10nm NMLs after 5 min or 20 min of isothermal annealing at 200 °C, 250 °C, 350 °C and 400 °C are shown in Fig. 3. Corresponding t2t XRD scans for the as-deposited and annealed AgCu8nm|AlN10nm NMLs specimens are presented in Fig. 4. No significant microstructural changes of the AgCu8nm|AlN10nm NMLs were detected by SEM/EDS after annealing at 200 °C and 250 °C. For these low annealing temperatures, only some adventitious carbon agglomerations and occasionally a very small Cu protrusion at the NML surfaces could be detected.

projection of the 3D NML, thus obscuring possible nanopores originating from incomplete coalescence of contacting nanograins.27-29

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Figure 3. Secondary electron image of the surface of the AgCu8nm|AlN10nm NMLs after 5 min or 20 min of isothermal annealing at (g,h) 200 °C, (e,f) 250 °C, (c,d) 350 °C and (a,b) 400 °C.

Figure 4. Theta-2theta (t2t) XRD scans of the AgCu8nm|AlN10nm in the as-deposited state, as well as after isothermal annealing for 5 mins at 200˚C, 250˚C, 350˚C and 400˚C. The asterisks indicate spurious radiation from a non-perfectly monochromatic source.



The corresponding XRD analysis did not detect any distinct reflections from a Cu-rich phase after annealing at 200 °C; i.e. only typical satellite peaks of the Ag(111), Ag(222), AlN(002) and AlN(004) reflections were observed, which originate from the periodic nanolaminated structure.15,19 The shoulder of the Ag(111) satellite peak in the t2t range of 40 - 45° can be attributed to Cu solute atoms in the supersaturated Ag[Cu] phase.5 A distinct Cu(111) reflection evolves from this shoulder of the Ag(111) peak upon heating at T ≥ 250 °C; the reflection is still relatively broad and weak at T = 250 °C, but becomes much sharper and more intense at T = 350 °C and T = 450 °C (see Fig. 4). This implies that the onset of phase separation becomes thermally activated at around 250 °C, but does not induce a noticeable degradation of the nanolaminated structure after 20 mins of annealing at T = 250 °C (see Figs. 3e and 3f). These experimental findings comply well with recent studies of the thermal stability of Ag-Cu40at.%/AlN and Al-Si12at.%/AlN NMLs by in-situ XRD, SEM and XPS, where first irreversible microstructural changes of the NMLs upon heating were always preceded by a phase separation of the nano-confined, supersaturated (i.e. metastable) alloy.5,18 For Al-Si12at.%/AlN NMLs, the phase separation proceeds by precipitation of solute Si atoms out of the Al[Si] fcc matrix in the T-range of 200 – 225°C.18 For Ag-Cu40at.%/AlN NMLs, the phase separation of the Ag[Cu] NLs into Cu and Ag domains occurs at a somewhat higher temperature of T ~ 250 °C. First significant microstructural changes of the Ag-Cu8nm|AlN10nm NMLs were observed by SEM/EDS after annealing at 350 °C and 400 °C (Figs. 3a-d). At the same time, the phaseseparated domains of Ag and Cu have coarsened, as reflected by sharper and more intense Cu(111) and Cu(222) reflections in the recorded XRD diffractograms. As evidenced by cross-sectional TEM,26 the progressive coarsening of the confined Ag and Cu domains can result in a domain width of tenths of nanometers (which by far exceeds the NL thickness), similar to the phase-separation process of Al and Si domains in AlSi12at.%/AlN NMLs.18 As evidenced from the planar SEM analysis, a widespread network of - 12 ACS Paragon Plus Environment

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line-shapes protrusions has appeared on the annealed NML surface (see Fig. 3a and, in particular, Fig. S2 in the supplementary material). These line-shaped networks have vastly extended after annealing at 400˚C, forming a hexagonal network with angles of roughly 120° between the contacting lines (see Fig. S2a). The line-shaped protrusions are adjoined by numerous dark patches (see Fig. 5a), which are distributed conterminous to the lines; their size and density roughly decreases with decreasing distance from the lines (cf. Fig. 5a, Figs. 3a,b and Fig. S2). The higher the line density, the higher the density of dark patches (see Fig. S2b). The interconnected line-shaped protrusions have an average width in the range of about 100 - 300 nm with a respective height in the range of about 30 – 50 nm, respectively. Strikingly, the width and height of "dead-ending" branches (i.e. line-shaped protrusions that are only connected to other lines at one outer end; cf. Fig. S2) decreases with increasing distance from the interconnection point to the tip of the branch, as accompanied by a decrease in the density and size of adjacent dark patches (see Fig. S2b).2 Comparable line-shaped networks have formed on the NML surfaces after 5 mins and 20 mins of annealing at 350 °C and 400 °C, which suggests that the formation of these patterns upon annealing occur rather abrupt. Higher magnification cross-sectional SEM micrographs of the AgCu8nm|AlN10nm NML after 5 min of annealing at 400 °C are presented with respective EDS elemental maps of Ag, Cu and O in Fig. 5. The SEM/EDS analysis shows that the line-shape protrusions are primarily constituted of Cu and O. The signal from Ag is strongly depleted at the locations of the protrusions; only rarely, localized smaller Ag droplets can be detected by SEM/EDS (see arrows in Fig. 5b). A SEM micrograph of a cross-sectional FIB cut perpendicular to a typical line-shaped protrusion is shown in Fig. 6.

Noteworthy, occasionally at the outer ending of such isolated branches, dark patches are no longer visible (cf. Fig. S2). 2



Figure 5 (a) Secondary electron micrograph and (b-d) EDS elemental maps of the AgCu8nm|AlN10nm NML after 5 min of isothermal annealing at 400 °C. The line-shaped protrusions form a hexagonal network with angles of 120° between the contacting lines. EDS mapping indicates that the line-shaped protrusions are primarily constituted of Cu and O; only rarely, localized smaller Ag droplets are detected (see arrows in Fig. 4b). The line-shaped protrusions are adjoined by numerous dark patches (see arrows in a), which correspond to voids in the uppermost Ag-Cu nanolayer(s) underneath the top AlN barrier layer.

Figure 6 (a) SEM micrograph of the surface of AgCu8nm|AlN10nm NML after 5 min of isothermal annealing at 400 °C, showing the formation of line-shapes Cu(O) protrusions and adjacent dark patches. (b) respective crosssectional SEM analysis perpendicular to the line-shaped protrusion of (a), evidencing that the dark patches correspond to voids in the uppermost Ag[Cu] nanolayers, as left behind by the mass transport of Cu, lateral along the internal interfaces, towards the line-shaped protrusion.


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It follows that the contiguous dark patches are voids underneath the top AlN barrier layer, which are concentrated in the uppermost Ag[Cu] nanolayers of the NML stack. These confined voids are left behind by mass transport of Cu parallel to the Ag[Cu]AlN interfaces towards the surface cracks, which have formed along preferential crystallographic directions in the textured AlN barrier layer(s). The hexagonal pattern of line-shaped protrusions originates from the in-plane texture of the six-fold symmetric Ag(111) and the hexagonally-packed AlN(0001) planes, as dictated by the orientation relationships (OR) between the α-Al2O3 substrate, the AlN barrier and the Ag[Cu] nanograins (see Figs. 2a and 2b).3 The video S4 in the supplementary material (see also caption text in Fig. S4) shows a series of SEM images obtained by sequential cross-sectional FIB slicing and imaging along a typical line-shape feature (referring to the AgCu8nm|AlN10nm NML after 5 min of annealing at 400 °C). The video S4 clearly evidences that the Cu, which has created the line-shaped Cu(O) protrusions, originates from the topmost alloy NLs. Occasionally, the surface cracks extend deeper into the NML and are covered by relatively large, approximately spherically-shaped Cu(O) protrusions with a typical diameter in the range of 0.5 - 1 µm. Fig. 7 shows a cross-sectional SEM micrograph of a FIB cut through such an extended surface crack and its associated Cu(O) protrusion. It follows that the surface crack has extended down two thirds through the NML system, thus allowing supply of Cu along the Cu/AlN and Ag/AlN interfaces from multiple phase-separated Ag-Cu NLs.

As verified in the present study, the hexagonal crack pattern can be modified into a rectangular crack pattern (with 4-fold-symmetry) by changing the α-Al2O3 substrate orientation from the (0001) prism plane to the (1102) basal plane. 3



Figure 7 (a) Cross-sectional STEM micrograph of a large Cu(O) protrusion, which has formed at the location of an extended surface crack in the AgCu8nm|AlN10nm NML after 5 min of isothermal annealing at 400 °C. The crack has extended down about two thirds through the NML stack. Corresponding EDS elemental mapping of (b) Ag, (c) Cu and (d) Al of the selected area in (a) show that the bright lines and larger white spots in (a) correspond to voids in the NML. The large black regions correspond to the accumulation of Ag.

The video S4 in the supplementary material shows a series of SEM images obtained by cross-sectional FIB slicing along the direction of the extended crack; it evidences that the Cu in the large protrusions originates from practically the entire NML stack. Large pores are detected just below the Cu(O) protrusions (corresponding to the white areas below the protrusion in Fig. 7). Furthermore, some Ag metal has collected in the region of the crack (corresponding to the large, dark regions in Fig. 7). The EDS mapping of the NML surface (Fig. 5) indicates that these Ag agglomerations are obscured by the large quantity of Cu(O) that lies on top of it. The voids left behind by the mass transport of Cu onto the NML surface are visible as bright horizontal lines in Fig. 7, which implies that the AlN barrier below these voids has not collapsed. However, some local deformation (bending) of the AlN barriers below the void structures can be observed. Effect of the AlN barrier thickness on directional outflow of Cu: SEM micrographs of the AgCu8nm|AlN8nm NMLs surface after 5 mins of isothermal annealing at 300 °C, 400 °C and 500 °C are shown in Fig. 8. A reduction of the AlN barrier thickness from 10 nm to 8 nm (for a constant Ag[Cu] NL thickness of 8 nm) results in a drastic decrease of the - 16 ACS Paragon Plus Environment

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density of line-shaped protrusions on the annealed NML surface at 400 °C: compare Figs. 3a and 8b. An isothermal annealing temperature as high as 500 °C is needed to induce more extensive fracturing of the AlN barrier layer with subsequent formation of line-shaped protrusions: see Fig. 8c. At some distance from these scarcely observed lineshaped protrusions, additional roughly spherical shaped protrusions with a diameter in the range of 10 – 200 nm (and occasionally even up to a dimeter of 800 nm) appear more or less randomly (i.e. homogenously distributed) on the annealed NML surface: see Fig 8f.

Figure 8. Secondary electron micrographs of the surface of the AgCu8nm|AlN8nm NML after 5 mins of isothermal annealing at (a, d) 300 °C, (b, e) 400 °C and (c, f) 500 °C. Panels (a, b, c) and (d, e, f) corresponds to high and low magnification, respectively.

A further reduction of the AlN barrier thickness from 8 nm to 4 nm (for a constant Ag[Cu] NL thickness of 8 nm) results in an increasing density of spherical protrusions and - 17 ACS Paragon Plus Environment


the absence of line-shaped protrusions on the NML surface, for the annealing at 400 °C and particularly at 500 °C: see Fig. 9. As evidenced by Figs. 9c-g, for the thinnest AlN barrier (thickness of 4 nm), the NML surface after annealing at 500 °C is densely and homogeneously covered with both submicron and nanometre sized droplets, while fracture-induced line-shaped protrusions are rarely (if at all) observed.

Figure 9. Secondary electron micrographs of the surface of the AgCu8nm|AlN4nm NML after 5 mins of isothermal annealing at (a, c) 300 °C, (b, d) 400 °C and (c, f, g) 500 °C. Panels (a, b, c) and (d, e, f) corresponds to high and low magnification, respectively. Panel (g) represent an even larger magnification of the NML surface of (f), evidencing a high density of submicron and nanometre sized droplets on the NML surface.

Cross-sectional STEM analysis of the annealed AgCu8nm|AlN4nm NMLs (5 mins at 500 °C) confirms the presence of a dense layer of nanoparticles on top of the outer AlN barrier layer: see Fig. 10a. EDS mapping of the STEM cross-sections indicates that these spherical protrusions are typically composed of single-phase Cu(O) (as for the lineshaped features observed for thicker AlN barrier layers), but frequently also consists of Ag (with possible small amounts of Cu in solid solution, as reflected by the Ag-Cu phase diagram at T = 400 °C): see Figs. 10b-d. Notably, the contact angle of the Cu protrusions on the AlN barrier layer is greater than that of the Ag droplets, indicating better - 18 ACS Paragon Plus Environment

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wettability of Cu on AlN (compare Figs. 10 b and c), which can be attributed to the formation of an interfacial CuO reaction layer.30 Indeed, the EDS mapping of the STEM cross-section detects a very thin oxide layer extending along the NML/air interface (see Figs. 10c,d).

Figure 10. (a) STEM micrographs with corresponding EDS Elemental mapping of (b) Ag, (c) Cu and (d) O for a cross section of the AgCu8nm|AlN4nm NML after 5 mins of isothermal annealing at 500 °C.

Cross-sectional TEM analysis confirms that the 4nm-thick AlN barrier layers after annealing are not fractured and still laterally continuous (with a relatively uniform thickness) (see Fig. 11b and footnote 1). Hence, a nanolaminated structure with an effective nano-confinement of the alloy is still ensured for AlN barrier thicknesses as small as 4 nm. Evidently, the 4-nm-thick AlN barriers behave much more ductile during heating, enabling local bending (instead of fracturing) of the AlN barrier layers. Consequently, a patterned outflow of Cu along networks of surface fractures is blocked; instead a dense and homogenous outflow of both Cu and Ag onto the NML surface occurs during heating of the AgCu8nm|AlN4nm NML. Strikingly, hardly any voids (as left behind by Cu and Ag migrated from the inner layers to the outer surface) were observed in the annealed AgCu8nm|AlN4nm NMLs, whereas such voids were omnipresent for the annealed AgCu8nm|AlN10nm (see Fig. 5a) and, to a lesser extend, also for the AgCu8nm|AlN8nm NML. For the annealed AgCu8nm|AlN4nm, any confined voids, as created - 19 ACS Paragon Plus Environment


by mass transport of Ag and Cu to the NML surface, have been closed by the deformation and partially sintering of the remaining AlN barrier layers: see Fig. 11.

Figure 11. TEM micrographs of a cross section of the AgCu8nm|AlN4nm NML after 5 mins of isothermal annealing at 500 °C. (a) A high density of CuO and Ag droplets is observed on top of the annealed NML, below which the AgCu NLs are partially emptied and AlN barrier layers sintered (no voids are observed). An enlargement of (b) an intact nanolaminated region of the NML, as well as (c) a region of sintered AlN barrier layers beneath the droplets in panel (a).

4.

Chemical driving forces and underlying mechanisms governing directional mass transport in Ag[Cu] /AlN nanomultilayers

First irreversible microstructural changes of the NMLs upon heating commence with a phase separation of the nano-confined, supersaturated (i.e. metastable) Ag[Cu] alloy into Cu and Ag domains occurs around T ~ 250 °C. The phase separation process is thermodynamically driven by the large positive enthalpy of mixing of Ag and Cu, - 20 ACS Paragon Plus Environment

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whereas the subsequent coarsening of the phase-separated Ag and Cu domains is mainly driven by the reduction of grain boundary density.5 The concurrent processes of phase segregation and domain coarsening, as accompanied by slight in-plane grain rotations of phase-separated Ag and Cu crystallites (resulting in an increasing in-plane coherency), all contribute to a lowering of the Gibbs energy of the system. The average (linear) thermal expansion coefficient (α) of the confined Ag[Cu] nanolayers (i.e. αAg  18×10-6 K-1 and αCu  17×10-6 K-1 and thus αAg[Cu]  αAg αCu) is roughly a factor 2-3 larger than those of the AlN barrier layers (αAlN  5×10-6 K-1) and the thick αAl2O3 substrate (αsapphire  8×10-6 K-1). Hence, in the absence of strain relaxation by plastic deformation,4 the confined Ag[Cu] NLs will, on average, experience a compressive bi-axial strain contribution upon heating, which is counteracted by tensile bi-axial strain contribution of the AlN NLs. These thermally-induced (in-plane) bi-axial stress contributions in Ag[Cu]/AlN NMLs can be relaxed by cracking of the top AlN barrier layer(s), as observed for the AgCu8nm|AlN10nm NML upon annealing at about T > 300 °C (see Fig. 3). Such fracturing of the top AlN barrier layer(s) in turn induces in-plane stress gradients in the NML from more compressive far away from the cracks to less compressive (or tensile) in the vicinity of the cracks. These in-plane stress gradients can drive directional mass transport of Cu along the semi-coherent Cu/AlN interfaces towards the surface cracks. Similar directional mass transport phenomena in thin-film systems under influence of in-plane (bi-axial) stress gradients have been reported for e.g. metal-induced layer exchange in Al/Si bilayer systems,31 Sn whisker growth of Sn coatings on Cu substrates32 and Al outflow in thermally-treated Al-Si12at.%/AlN NML systems.18 The relaxation of compressive stresses within the phase-separated Ag-Cu NLs is also evidenced by a corresponding peak shift of the Ag(222) reflection to lower thetaThe isothermal annealing treatment is designed to be extremely rapid (see Sec. 2) and therefore gives very little time for the substrate/NML system to relax any thermal stress during fast heating (and cooling). It may thus be assumed that any thermal stresses during fast heating cannot be relaxed. 4



2theta values after the annealing at 350 °C and 400 °C (see Fig. 4). Noteworthy, the thermally-induced (in-plane) bi-axial stress contributions in Ag[Cu]/AlN NMLs can also be relaxed by local bending (hillock formation) of the top barrier layers, which dominates for the thinnest AlN barrier layers (see Fig. 11). Annealing of the AgCu8nm|AlN10nm NML invokes preferential mass transport of Cu towards the surface; i.e. mass transport of Ag to the surface cracks is negligible (see Fig. 5). This experimental finding is rather striking, since the slightly higher homologous temperature of Ag as compared to Cu (i.e. T/Tm [K], which equals 0.54 for Ag and 0.50 for Cu at T = 400 °C, respectively) may suggest a slightly higher atomic mobility of Ag as compared to Cu. Alternatively, the self-diffusion coefficients of Ag and Cu at T = 400 °C can be taken as a very rough measure for comparing the atomic mobilities of Ag and Cu: the self-diffusion coefficients of Ag and Cu at T = 400 °C equal 3×10-19 cm2/s and 6×10-21 cm2/s, respectively.33 This implies that the self-diffusion coefficient of Ag at T = 400 °C is a factor 200× higher than the respective self-diffusion coefficient of Cu. Hence it may be naively argued and expected that mass transport in the confined AlN/Ag[Cu] is preferentially realized by the diffusion of Ag. Strikingly, the opposite is observed: i.e. mass transport upon annealing preferentially occurs by the outward diffusion of Cu. The much higher atomic mobility of Cu (as compared to Ag) in the studied Ag[Cu]/AlN NMLs can be partly rationalized on the basis of the significantly higher tracer diffusion coefficients for Cu in Ag (  5.3×10-15 cm2/s at T = 400 °C) as compared to Ag in Cu ( 4.7×10-16 cm2/s at T = 400 °C).33 Moreover, as postulated based on DFT predictions of the defect structures of the semi-coherent Ag/AlN and Cu/AlN interfaces, the atomic mobility of Cu along the disordered Cu/AlN interface may be higher as compared to the atomic mobility of Ag along the more coherent Ag/AlN interface.5 As such, the directional mass transport of Cu along Cu/AlN interfaces will also be favoured. Finally, Cu has a much higher affinity for oxygen than Ag, which drives O-induced chemical - 22 ACS Paragon Plus Environment

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segregation of Cu to the NML surface.5 The chemical-segregation is driven by the chemical potential gradients of Cu and O,34-35 as established both parallel and perpendicular to the surface cracks. Namely, the surface crack can be easily penetrated by oxygen,15 thereby creating localized regions of high O activity in the vicinity of the cracks, which drives chemical segregation of Cu along the Ag/AlN and Cu/AlN interfaces to these O-enriched regions.34-35 This also explains the partial oxidation of the lineshaped Cu protrusions. The pattern of line-shaped Cu protrusions on the AgCu8nm|AlN10nm NML surface remains more or less unchanged with increasing annealing time (at constant annealing temperature). This implies that the surface fracturing of the AlN barrier layers with subsequent directional mass transport of Cu along the internal interfaces towards the surface cracks occurs at a very fast rate and is already completed within a relatively short time interval of a few minutes, in accord with previous investigations of the thermal stability of Ag/AlN and AgCu/AlN NMLs by real-time HT-XRD at the synchrotron.5,15 Strikingly, the diffusion length of Cu along the Cu/AlN interfaces to the cracks under the influence of the acting in-plane gradients in the residual stress and the chemical potential (see above) extends to several micrometers (e.g. from isolated dark patches to the nearest line-shape feature; cf. see Figs. 3a,b and 5), which by far exceeds the typical solid-state diffusion lengths in binary alloys at comparably-low temperatures (of, say, several tenths to maximum a few hundreds of nanometers). As previously emphasized, the AlN nano-grains (as nucleated on the Ag[Cu] surface at the onset of each AlN deposition step) have not yet completely coalesced into a continuous and compact AlN nanolayers.27-29 Consequently, the 4-nm-thick AlN barrier layers still have a relatively open grain boundary structure (resulting in tiny nanopores at the AgCu8nm|AlN4nm NML surface; see Fig. S1b), which provides alternative pathways for fast diffusion of Ag and Cu to the NML surface during heating. Consequently, for AlN - 23 ACS Paragon Plus Environment


barrier thicknesses as small as 4 nm, thermally-induced (in-plane) bi-axial stress contributions in Ag[Cu]/AlN NMLs can be easily relaxed by homogenous outflow of Cu (and Ag) through the relatively open grain boundary structure of the AlN barrier layers. The quantity of local outflow of confined metal will differ depending on the local degree of coalescence of the AlN nanograins, which rationalizes the observed spread in the size of the spherically shaped metal protrusions. Any voids, as created by mass transport of Ag and Cu to the NML surface, have been largely annihilated by the deformation and partially sintering of the relatively ductile 4nm-thick AlN barrier layers (see Fig. 11), as driven by the associated reduction of surface and interface energies.17 In summary, the reported findings convincingly demonstrate that the AlN barrier thickness has a crucial effect on the mechanism of directional mass transport of the nano-confined alloy to the NML surface upon heating. The implementation of relatively thick (i.e. 10-nm-thick) AlN barrier layers, on the one hand, invokes patterned fracturing of the AlN barrier layers upon heating with subsequent fast mass transport of Cu along internal interfaces to the surface cracks, resulting in the formation of an interconnected network of Cu(O) lines on the annealed NML surface. Much thinner (i.e. 4-nm-thick) AlN barrier layers, on the other hand, act as nanoporous membranes upon heating, enabling direct migration of Cu and Cu through the AlN barrier layer without fracturing, resulting in the formation of a dense and homogenous distribution of CuO and Ag droplets on the annealed NML surface (cf. Figs. 9b-f). These findings might open new routes for benign bonding of micro- and nano-scaled systems at ever-lower temperatures. Previous thermodynamic predictions17 approve our assertions that the concurrent processes of (i) phase separation of the nano-confined Ag[Cu] alloy, (ii) subsequent coarsening of the phase-separated Ag and Cu domains, (iii) thermal stress relaxation and (iv) fast outflow of metal to the NML surface (and its partial oxidation), all contribute to a lowering of the Gibbs energy of the NML system and thus provide the effective - 24 ACS Paragon Plus Environment

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thermodynamic driving force for the transformation of NML structure upon heating. As evidenced in the present study, the thermodynamics and kinetics of this structural transformation critically depend on the complex interplay between the different NML design parameters, such as the thickness, composition, grain size, texture, (defect) structure, and residual stress state of the alloy and barrier layers, as well as the defect structure of the respective alloy/barrier interfaces. Our current research will continue to explore and optimize the complex interplay between these different NML design parameters, while targeting at directional mass transport at even lower temperatures (i.e. T < 300 °C).

5.

CONCLUSIONS

The phase stability and atomic mobility of nano-confined Ag-Cu40at% alloys in Ag[Cu]/AlN NMLs upon isothermal annealing in air were studied. In particular, specific knowledge is obtained on how to tune the AlN barrier layers for achieving fast and directional mass transport of the confined metal to the NML surface upon heating. In the as-deposited state, the nano-confined AgCu alloy nanolayers (NLs) are composed of a fcc matrix of Ag nano-grains, which are supersaturated by Cu (further denoted as Ag[Cu]). The nano-crystalline Ag[Cu] and AlN NLs are of uniform thickness and possess a strong in-plane and out-of-plane texture, which originates from the orientation relationship between the -Al2O3(0001) substrate, the AlN barrier and the Ag[Cu] NL according to Ag{111}AlN{0001}Al2O3{0001}. The coalescence of individual AlN islands into a compact AlN barrier is not yet fully completed for barrier thicknesses as small as 4 nm, resulting in relatively open grain boundary structure. The lower degree of coalescence of the thinnest AlN barriers is also reflected by a less pronounced in-plane coherency of the Ag[Cu]AlN interface.



For the annealed NMLs, first significant microstructural changes were observed after annealing at T  250 °C. The AgCu8nm|AlN10nm NML (and to a much lesser extend the AgCu8nm|AlN8nm NML) develops an interconnected network of fractures after the annealing at T > 300 °C, which is completely decorated with Cu(O) protrusions. The fracturing of the AlN barrier layers relaxes the thermally-induced in-plane bi-axial stresses in the AgCu/AlN NML and invokes very fast, stress- and chemically driven mass transport of Cu along internal interfaces to the cracked regions, forming Cu(O) lineshaped protrusion on the annealed NML surface. The directional mass transport process leads to the creation of confined voids within the NML adjacent to the line-shape protrusions, which are primarily concentrated in the topmost alloy NLs. For the annealed AgCu8nm|AlN4nm NML (corresponding to the thinnest AlN barrier layer of 4 nm), such fracturing of the AlN barriers layer with associated formation of Cu(O) line-shaped protrusions and confined voids is not observed. Instead, thermally induced stresses are relaxed by outflow of Cu (and Ag) through the relatively open grain boundary structure to the NML surface at T > 300 °C. Any confined voids, as created by mass transport of Ag and Cu to the NML surface, are annihilated by bending and partially sintering of the relatively ductile AlN barrier layers.

ASSOCIATED CONTENT Supporting information: The Supporting Information is available free of charge on the ACS Publications website. - AFM analysis of AgCu8nm|AlN4nm and AgCu8nm|AlN10nm NML surfaces (S1) - Photograph of the experimental setup for isothermal annealing in air (S2) - SEM analysis of the AgCu8nm|AlN10nm NML surface after 5 min of annealing at 400 °C (S3)


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- Cross-sectional FIB/SEM imaging along the direction of a typical line-shape feature (for AgCu8nm|AlN10nm NMLs after 5 min of annealing at 400 °C (Video S4) - Cross-sectional FIB/SEM imaging through an extended crack and its associated relatively large Cu(O) protrusion on top (for AgCu8nm|AlN10nm NMLs after 5 min of annealing at 400 °C (Video S5).

AUTHOR INFORMATION Corresponding Author * Dr. Lars P.H. Jeurgens, Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Joining Technologies & Corrosion, Ueberlandstrasse 129 8600 Duebendorf, Switzerland. Tel +41 58 765 4053. Email: [email protected] ORCID Claudia Cancellieri: 0000-0003-4124-4362 Mirco Chiodi: 0000-0003-0418-0785 Jolanta Janczak-Rusch: 0000-0002-9161-5822 Lars P.H. Jeurgens: 0000-0002-0264-9220 Present Addresses † The School of Engineering and Materials Science, Queen Mary University of London, Mile End Road, London E1 4NS (email: [email protected]) Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.



Funding Sources We acknowledge the support received from EU FP7-PEOPLE-2013-IRSES Project EXMONAN-Experimental investigation and modeling of nanoscale solid state reactions with high technological impact (Grant no. 612552).

ACKNOWLEDGMENT We are grateful to Dr Hans-Ruedi Elsener for designing and performing the isothermal annealing experiments and to Dr. Luchan Lin for the AFM analysis of the as-deposited NML surfaces. The electron microscopy center of Empa, as well as the Empa laboratory for Transport at Nanoscale Interfaces, are acknowledged for providing access to the FIB, SEM and TEM facilities.

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