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DOI: 10.1002/adem.201500578
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Atomic-Scale Structural and Chemical Study ofColumnar and Multilayer Re–Ni ElectrodepositedThermal Barrier Coating**

By Sung-Il Baik, Alla Duhin, Patrick J. Phillips, Robert F. Klie, Eliezer Gileadi,David N. Seidman and Noam Eliaz*

Rhenium alloys exhibit a unique combination of chemical, physical, and mechanical properties thatmakes them attractive for a variety of applications. Herein, we present atomic-scale structural andatomic part-per-million level three-dimensional (3D) chemical characterization of a Re–Ni coating,combining aberration-corrected scanning transmission electron microscopy (STEM) and atom-probetomography (APT). A unique combination of a columnar and multilayer structure is formed by single-bath dc-electroplating and is reported here for the first time. Alternating thicker Re-rich and thinnerNi-rich layers support a mechanism in which Ni acts as a reducing agent. The multilayers exhibithetero-epitaxial growth resulting in high residual shear stresses that lead to formation of corrugatedinterfaces and an outer layer with mud-cracks.

1. Introduction high cost. Interest in electroless deposition andelectrodeposition[5–16] of Re and its alloys has recently

Rhenium (Re) is a refractory metal with a uniquecombination of chemical, physical, and mechanical propertiesthatmakes it attractive for high-temperature, catalytic, energy,electrical, biomedical, and other applications, in spite of its

[*] Prof. N. Eliaz, A. DuhinDepartment of Materials Science and Engineering, Tel-AvivUniversity, Ramat Aviv, Tel Aviv 6997801, IsraelE-mail: [email protected]. S.-I. Baik, Prof. D. N. SeidmanDepartment of Materials Science and Engineering, Northwest-ern University, Evanston, Illinois 60208, USANorthwestern University Center for Atom Probe Tomography(NUCAPT), Evanston, Illinois 60208, USADr. P. J. Phillips, Prof. R. F. KlieDepartment of Physics, University of Illinois at Chicago,Chicago, Illinois 60607, USAProf. E. GileadiSchool of Chemistry, Tel-Aviv University, Ramat Aviv, Tel Aviv6997801, Israel

[**] N. Eliaz thanks the McCormick School of Engineering andApplied Science (MEAS) at Northwestern University (NU) aswell as the International Institute for Nanotechnology of NUand Tel-Aviv University (TAU) for awarding him an Eshbachscholarship and a travel grant, respectively. This research wasperformed in part with financial support from the US Air ForceOffice of Scientific Research (AFOSR, grant number FA9550-

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[1,2] [3,4]

emerged. Electroplating of pure Re yields poor depositionresults. The addition of salts of the iron-group metals (Ni, Fe,and Co) to the plating bath improves, however, the coatingdramatically, allowing deposition of thick layers of the

10-1-0520) and the Israel Department of Defense (grant number4440258441). The authors thank M. Levenshtein andM. Cohen-Sagiv of the Biomaterials and Corrosion Lab atTAU for their technical help. Atom-probe tomography measure-ments were performed at the Northwestern University Centerfor Atom-Probe Tomography (NUCAPT), and the LEAPtomograph was purchased and upgraded with funding fromthe NSF-MRI (Grant No. DMR 0420532) and ONR-DURIP(Grant Nos. N00014-0400798, N00014-0610539, N00014-0910781) programs. NUCAPT is a shared facility of theMaterials Research Center of NU, supported by the NationalScience Foundation’s MRSEC program (Grant No. DMR-1121262). We are also grateful to the Initiative for Sustainabil-ity and Energy at Northwestern (ISEN) for grants to upgradeNUCAPT’s capabilities. The acquisition of the JEOL JEM-ARM200CF at the University of Illinois at Chicago (UIC) issupported by an MRI-R2 grant from the National ScienceFoundation (Grant No. DMR-0959470). Support from the UICResearch Resources Center is also gratefully acknowledged.(Supporting Information is available online from Wiley OnlineLibrary or from the author).

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S.-I. Baik et al./Atomic-Scale Structural and Chemical Study of Re-Ni. . .

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corresponding alloys with very high Re-concentrations (up to
93 at%) and high Faradaic efficiency (FE, up to 96%).[3,5] Amechanism by which the presence of the iron-group metal(Me) increases the rate of Re deposition is proposed.[6,7] Itinvolves the electro-reduction of Me2þ cations on thesubstrate’s surface, followed by chemical reduction of theperrhenate (ReO�

4 ) ion, probably to the less stable rhenate ion(ReO�

3 ), which is then reduced electrochemically to metallicRe, forming a Re-Me alloy.

Electrocrystallization is influenced by surface preparationof the substrate and surface defects, bath chemistry, oxygenlevel, pH, stirring, the applied potential (or current density),temperature, cell geometry, etc.[17] The chemical compositionand microstructure of the electrodeposits have a significanteffect on their mechanical and corrosive (both aqueous andoxidation at high temperatures) properties. To gain an in-depth understanding of the electrochemical process, wepresent herein the first atomic-scale structural and chemicalcharacterization of a Re–Ni coating. The correlative study[18–20]

combined picosecond ultraviolet laser-assisted local-electrodeatom-probe (LEAP) tomography[21–23] and spherical aberra-tion-corrected scanning transmission electron microscopy(STEM)[24,25] performed on the same needle-like sample thathad been prepared by the focused ion-beam (FIB) microscopylift-outmethod.[26,27] Thus, insightson theelectrocrystallizationofRe–Ni are provided byboth structural characterization at theatomic scale and atomic part-per-million (at. ppm) 3D chemicalmapping, for the first time. We note that, while field-ionmicroscopy (FIM) and LEAP tomography have already beenused for characterization of radiation damage, grain boundary(GB) segregation and multiple-step formation at elevatedtemperatures in pure bulk Re,[28] as well as for analyses of thedistribution of Re in Ni-base superalloys,[29,30] radiation-induced precipitation in W–Re alloys,[31] and Re-monolayerdeposition on a W surface,[32] this article presents the firstatomic-scale tomographic study of Re-based coatings.

2. Experimental Section

2.1. Material PreparationThe substrate was 316L stainless steel sheet. Surface

preparation included etching in a 1:1 HCl:HNO3 solution for1min, followed by activation at room temperature in aPd-citrate buffer solution (pH 2.5) for 2min.[3,4] Electro-plating was performed in a bath consisting of 22.67mmol L�1

ammonium perrhenate (NH4ReO4), 34mmol L�1 citric acid(complexing agent), and 34mmol L�1 of nickel sulfamate(Ni(NH2SO3)2). All components were dissolved in deionizedwater (Direct-Q3 from Millipore). The exposed surface areaof the cathode was 1.08 cm2. Two platinum sheets were usedas the anode (counter electrode), and a saturated Ag/AgCl/KCl as the reference electrode. Additional information on theexperimental procedure is provided elsewhere.[5,6] The pHwas adjusted to 2.3–2.4, and a thermostatic bath (MRC B300)was used to maintain the temperature at 60 �C. A currentdensity of 50mA cm�2 was applied for 1 h, using a

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potentiostat/glavanostat (Princeton Applied Researchmodel 263A).

2.2. Preliminary Characterization of the ElectrodepositThe Faradaic efficiency (FE) was calculated from the mass

gained, the charge passed, and the chemical composition ofthe deposit. The partial deposition current densities werecalculated from the mass gained and the chemical composi-tion of the deposit.[9] The latter was determined by energy-dispersive spectroscopy (EDS), using a liquid-nitrogen-cooledOxford Si EDS detector attached to an environmentalscanning electron microscope (ESEM, Quanta 200 FEG fromFEI). The chemical composition was calculated as the averageof five measurements at different spatial positions. The ESEMwas also used to characterize the surface morphology of thecoating. The thickness of the coating was estimated based onmass gain measurements, taking into account the atomicfraction of each element in the alloy.

2.3. Sample Preparation for APT and STEMThe Re–Ni film on a 316L stainless steel was covered by an

�100mm thick Ni over-layer to protect the top surface,utilizing an Arþ ion-beam sputter/etching system (IBS/e)operating at 8 kV and 6mA. Samples for both atom-probetomography (APT) and scanning transmission electronmicroscopy (STEM) characterization were prepared by thelift-out method, utilizing a dual-beam FIB microscope (HeliosNanolab, FEI Co., Hillsboro, OR, USA).[33,34]

2.4. Aberration-Corrected STEM CharacterizationHigh-resolution microscopic, crystallographic, and chemi-

cal analyses were performed using the probe aberration-corrected JEOL JEM-ARM200CF at the University of Illinois atChicago, equippedwith a 200 kV cold-field emission gun, anda high-angle annular dark-field (HAADF) detector. A 28mradprobe convergence semi-angle was used for HAADF(Z-contrast) imaging, with an inner detector angle of 110mrad. An Oxford X-max 80 SDD X-ray detector was used forEDS analysis. Complementary STEM characterization wasperformed at Northwestern University, using a JEOL 2100FTEM. The selected area electron diffraction (SAED) patternswere indexed using commercial software programs, Crystal-Maker and Single Crystal.

2.5. Atom-Probe Tomography (APT) AnalysesAn ultrahigh vacuum (gauge pressure less than 2� 10�11

Torr) local-electrode atom-probe (LEAP) tomograph[21–23]

(LEAP 4000X-Si) was utilized at Northwestern University tomeasure the chemical compositions of APT nanotip samples.Picosecond pulses of ultraviolet laser light (wavelength¼ 355nm) were applied to evaporate individual atoms from APTnanotips at a pulse repetition rate of 100kHz, a laser energy of15–30 pJ pulse�1, an average detection rate of 0.01 ions pulse�1,and a nanotip temperature of 60.0� 0.3K. The nanotips weremaintainedat apositivepotential,while the evaporationof ionswas triggered by the picosecondUV laser pulses. The times-of-

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flight of the detected ions were utilized to producethree-dimensional (3D) reconstructions using theprogramIVAS,version3.6.6.Thecompositionsof theRe–Niphasesweremeasuredbytime-of-flightmass-to-charge state ratio spectra (Supporting Informa-tion S1).

Fig. 1. SEM images of the Re–Ni electrodeposit on a 316L stainless steel substrate. (a, b) Top surfaceview revealing a mesoscale colony structure at (a) low- and (b) high-magnification. Intercolonyboundaries, crevices, and 3D nanometric islands formed by a Volmer–Weber growth mechanism areevident. (c) Dual-beam FIB microscope Gaþ-ion-milled cross-sectional view of the rectangular area in(a). The coating consists of a bottom single-phase layer, I, and a top multilayer zone, II, see inset forclarity. (d) Fracture surface revealing a columnar grain structure with several zones with different grainsizes across the coating. Layers with small grains are observed on the substrate and at the mid-thickness.The outer layer consists of a nodular structure. In between these layers, larger columnar grainsdominate.

3. Results

3.1. Electroplating of Re–NiMeasurements of the mass gain, electric charge

passed during electrodeposition, and the chemicalcomposition of the coating enabled us to calculate:Faradaic efficiency (FE)¼ 21.3%;partial currentden-sities iRe¼ 4.302mAcm�2 and iNi¼ 0.916mAcm�2;and a coating thickness of about 6mm. Scanningelectron microscopy (SEM)-energy dispersiveX-ray spectroscopy (EDS) analyses indicated thefollowing chemical composition: 57.3� 1.8 at%Re and 42.7� 1.8 at% Ni, Table 1. In this study, alower value of the FE was obtained comparedto our prior studies.[5–7,9] The reason for thisdifference is most likely the lower pH in the presentstudy (2.3–2.4 vs. 5.0), for which visually nicercoatings were obtained with a concomitant en-hancement of the hydrogen evolution reaction.

3.2. SEM Characterization of Surface Morphologyand Cross-Section

SEM secondary electron (SE) images of the topsurface of the coating are displayed in Figure 1a,b.
Amesoscale colony surface structure (also known as a nodularor cauliflower structure) is evident, with clear intercolonyboundaries and surface grooves. The typical colony size is11.2� 2.0mm; each colony consists of hundreds of small grainsthat are �45nm in diam. The appearance of these 3D islandsindicates a Volmer–Weber (VW) growth mode.[35,36] The mostimportant parameters determining the mode of growth of asubstance on a substrate made of another material are thedeposit/substratebindingenergyand the latticeparameterandstructuralmisfit between them.For a small supersaturationandnegligible kinetic influences, a depositmaygrowon a substrate
Table 1. Comparison between the chemical compositions (at%) determined bySEM-EDS (top surface), STEM-EDS (distinct layers), and LEAP tomography(inter-columnar grain boundary).

Re[at%]

Ni[at%]

O[at%]

H[at%]

SEM-EDS 57.3 42.7 x xSTEM-EDS (bright layer) 58.6 41.4 x xSTEM-EDS (dark layer) 49.5 50.5 x xLEAP (bright layer withincolumnar grain)

61.3 37.6 0.7 0.4

LEAP (at the inter-columnar GB) 42.3 52.7 3.8 1.2

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via different modes. The VW mode is preferred, regardless ofthe lattice parameter and structural misfit, when the bindingenergy between adsorbed deposited ions is greater than thebinding energy between an adsorbed ion and a substrate atom.The formation of nodules (colonies) may be the outcome ofnucleation extending laterally fromanexistingstructural defect(e.g., dislocation) and sweeping impurity atoms (e.g., hydro-gen) ahead of the growth front. The impurity atoms would,therefore,bemoreconcentratedwherenodulesmeetandinhibitoutgrowth, i.e., perpendicular to the surface.[37] Schuh et al.suggested that the colony structure may involve bothcrystallographically related groups of grains and subtlevariations in composition at the colonies’s boundaries.[38]

Surface crevices (or cracking) in electrodeposits may berelated to decomposition of unstable hydrides during or soonafter plating, which involves significant volumetric contrac-tions and internal stresses.[37,39] Examples include decompo-sition of hexagonal nickel hydride Ni2H to pure hexagonal orface-centered cubic (fcc) Ni,[9] or decomposition of hexagonalchromiumhydride to body-centered cubic (bcc) chromium.[39]

Even in the absence of hydrides, co-deposition of hydrogenmay result in intrinsic stresses, either tensile (if hydrogendiffuses from the surface layer to the layers beneath it) orcompressive (if hydrogen remains in the surface layer and fillsvoids).[37] Such macroscopic intrinsic stresses are likely to be

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larger in plating systems with a low FE. Residual stresses may
also be related to the manner in which metal grains crystallize(e.g., variation in grain size), or to defects incorporated in thecoating (e.g., microscopic voids, metallic impurities, disloca-tions, and interfaces).[40] Macroscopic extrinsic stresses aremainly due to epitaxial growth and a significant differencebetween the interatomic spacing of the substrate and depositmaterials. When the deposit thickness reaches a critical value,misfit dislocations start developing, thereby relieving thesestresses. Macroscopic intrinsic stresses develop within thedeposit when it is restrained from shrinking or expanding. Forexample, when 3D epitaxial crystallites coalesce before thespaces between them have been completely filled, the surfacelayer of a deposit tries to shrink. The restraint by the layersunder the surface causes tensile stresses, which become largeras the grains become smaller.[37]

Abermann classified the time (thickness) dependent stressbuild-up patterns in VW vapor deposited thin films into twotypes.[41] Type I, which is relevant to this study, is typical ofhigh melting point elements, e.g., Re, W, Ta, etc. For theseelements, the final microstructure often has high aspect ratiocolumnar grain structure. Tensile stresses are developedduring growth, but when growth stops, the final film’s stressstate usually does not change, or changes only slightly, withincreasing time. The stress build-up in electrodeposited filmsis, however, significantly different from that occurring invapor-deposited films.

Figure 1c is a SEM image of a Gaþ ion-milled cross-sectionof the Re–Ni coating. The thickness of the deposit is10.6� 0.6mm, and it consists of two distinct zones: i) abottom single-phase zone (2–3.5mm thick, indicated by I); andii) a top multilayer zone (indicated by II). Based on onlycontrast, the chemical composition within zone I is nearlyuniform. In distinction, compositionally modulated multi-layers (often termed banded, striated, or lamellar structures)are evident in zone II (note the alternating bright and darklayers). The interlamellar spacing (λ) is �200–300 nm. Theinset in Figure 1c demonstrates that the layers in zone II have awavy or corrugated morphology. The darker region adjacentto the outer surface consists of nanograins that are arranged incolonies.

Figure 1c does not display any evidence of columnargrains. During surface preparation using the dual-beam FIBmicroscope, it became clear that milling with Gaþ ionssmooths the exposed surface, masking some microscopicfeatures, and revealing crack propagation (note the crackingin Figure 1c). It is evident that the surface crevices (cracking)observed in Figure 1a and b do not penetrate through thecoating but are limited to the top surface layers. Whenobserving freshly fractured surfaces, a columnar structurewas exposed, Figure 1d. A nodular structure with nanograinsis evident near the top surface. The thickness of the zone withthis structure is �1.2mm. These groups of nanograins mergeinto large columnar grains with a similar crystallographictexture. Small grains are also present at the bottom portion ofthe coating, within �500 nm from the interface between the


coating and substrate. These grains reflect numerousprecipitation sites (instantaneous nucleation) of the initialdeposited layer on the stainless steel substrate. A layerconsisting of fine columnar grains is also observed atapproximately mid-thickness. Layers consisting of largecolumnar grains exist between this center layer and the topand bottom layers with a fine grain structure (i.e., a few tensof nm in diameter).

3.3. TEM and Cs-Corrected STEM Microstructural andCrystallographic Characterization

For high-resolution microstructural and crystallographiccharacterization, TEM and aberration-corrected STEM wereutilized. Figure 2a displays large columnar grains from thezone beneath the nodular structure seen in Figure 1d. It isapparent that each columnar grain consists of alternatingthick (bright) and thin (dark) layers. STEMhigh-angle annulardark-field (HAADF) imaging has strong compositionalsensitivity and provides atomic number (Z) contrast.[24] Thus,the bright layers in Figure 2a are associated with a phase richin Re (Z¼ 75), whereas the dark layers are associated with aphase rich in Ni (Z¼ 28). The Ni-rich phase represents avolume fraction of 0.150� 0.006 of the multilayer zone. Acrack along the intercolumnar GB is apparent. As evidentfrom Figure 1c, the intercolumnar GBs are the preferred sitesfor crack propagation. The Ni-rich layers in Figure 2a arecurved, with their convex surface in the direction of the topouter surface of the coating. Additionally, each layer isdiscontinuous when crossing the intercolumnar GB becausesome shift appears. The curvature and discontinuity of thesethin layers, along with the tendency for crack initiation andpropagation along the intercolony and intercolumnar grainboundaries, indicate the existence of high internal stresses.

A high-magnification image of the interfaces between aNi-rich layer (dark) and the adjacent Re-rich layers (bright) isdisplayed in Figure 2b.While the lower interface of theNi-richlayer in this figure is smooth, the upper interface is very rough(Dz¼ 11.8� 3.5 nm, compared to its thickness of 18.2 nm) andexhibits a wedge shape. This roughness is probably related tohetero-epitaxial growth of the two phases. Dislocation arraysare also apparent within the Ni-rich layer (indicated byarrows), some of which extend into the Re-rich phase at theapex of the wedge. The ratio between the numbers ofdislocations in the two phases is nearly 3:1 for the Ni- andRe-rich phases, respectively, because Ni is softer than Re andcan sustain more plastic deformation. The difference in theinterplanar spacing of the two phases results in repeatedexpansions and contractions within the multilayer coating.Dislocations release the elastic strain energy because the extraatomic planes behave as small wedges, speaking qualitatively.The dislocations can also compensate for the misorientationbetween colonies at the top surface. The coalescence of twolow-angle grains results in one grain boundary with a smallergrain boundary energy. The result of strains and dislocationarrays in the multilayer structure is a strongly corrugatedinterface.


Fig. 2. STEM-HAADF imaging and SAED analyses. (a) Alternating layers of a Re-rich phase (bright) and a Ni-rich phase (dark) within large columnar grains. A crackemanating along an inter-columnar GB, curvature of the Ni-rich layers, and a relativeshift of these layers on both sides of an inter-columnar GB, are observed. (b) High-magnification image of the two interfaces between a Ni-rich layer (dark) and adjacentRe-rich layers (bright). The upper interface in this image is rough and a high density ofdislocations is apparent on this side. These two features are related to the epitaxialgrowth of the two phases and to stress relaxation at the interface, respectively. (c) SAEDpattern displaying the crystallographic orientation at the interface. The diffractionpattern from an hcp Re-based phase is delineated by green solid-lines and exhibits an1�101��

orientation, as in the inset. The diffraction pattern from a fcc Ni-based phase isdelineated by red dotted lines and it exhibits an 010½ � orientation.



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By analyzing the structure of many layers we find that inmost cases two adjacent layers within a columnar grainexhibit hetero-epitaxial growth, with hexagonal close-packed(hcp) rhenium (lattice parameters, a¼ 0.2761 nm, c¼ 0.4456nm[42]). The selected area electron diffraction (SAED) patternin the inset of Figure 2c exhibits a 1�101��

zone-axis orientationof the two columnar grains encircled in Figure 2a. A smallrotation exists, �3�, as indicated by the two arrows, whichreveals the misorientation between the two columnar grains.The coalescence of two small-angle grain boundaries results inone grain boundary with an excess energy that is smaller thanthat of the initial two grain boundaries. On the basis of theSAED analysis, the formation of the metastable hcp-Ni phaseexhibits hetero-epitaxial growth on hcp Re-rich layers in mostof the columnar grains. By analyzing carefully the relativeorientation of the two Re-rich layers on both sides of a Ni-richlayer it is shown that some columnar grains have a change ofcrystallographic orientation on the Ni-rich layer. It is evidentthat on one side the boundary is coherent or semi-coherent,while on the other side the boundary is incoherent with adifferent orientation. This is the reason for the one-sidedcorrugations seen in Figure 2b. The Ni-rich layer inFigure 2b has a straight interface with the bottom Re-richlayer because it has another crystallographic orientation andhence epitaxial strains are not involved. On the other side,when the Ni-rich layer is sufficiently thick (�20 nm in thiscase) and rhenium is deposited on top of it, the interface iscorrugated due to hetero-epitaxial effects and because Ni issofter than Re. Epitaxial growth is demonstrated by the SAEDpattern, Figure 2c, acquired from the areamarked by a circle inFigure 2b. The diffraction pattern of the hcp-Re phase isindicated using solid green-lines and it exhibits the same1�101�hcph

orientation as in the inset, while the diffractionpattern of the fcc-Ni phase is marked with red dotted-linesand it exhibits a 001½ �fcc orientation. The measured latticeparameter of this fcc-Ni-rich phase (a¼ 0.384 nm) indicatesthat a significant amount of Re is dissolved in it.

The local atomic arrangement and crystallographic struc-ture are revealed utilizing the aberration-corrected STEM-HAADF images, their fast Fourier transform (FFT), and theirsimulated images; all of which are displayed in Figure 3.Figure 3a–c displays the fcc-Ni structure with a [001]orientation, the hcp-Re structure with a 1�101��

orientation,and the hcp-Re structure with a 2�423��

orientation. Thepositions at which the high-resolution images were acquiredare marked in Figure 2b. The two phases in Figure 3a,b sharetheir (010)fcc and 01�11Þhcp

�stacking planes with interplanar

spacings of 0.192 and 0.211 nm, respectively. This �9%difference is related to the strain state of this multilayercoating. This observation is in agreement with the SAEDpatterns, Figure 2c. Figure 3c displays a different orientationof the Re-rich phase, below theNi-rich phase in Figure 2b. Theplanes perpendicular to the substrate are close to 01�11Þ�

and10�10Þ�

in the hcp-Re-rich phase, and to (010) in the fcc-Ni-richphase. The corresponding planes are not close-packed in thehcp and fcc structures. This observation agrees with the


Fig. 4. Re and Ni STEM-EDS profiles along the line indicated in Figure 2a.The Re-concentration in the bright layers is higher than in the dark layers, in agreementwith the STEM-HAADF Z-contrast image in Figure 2a.

Fig. 3. Aberration-corrected STEM-HAADF images acquired at three different locations,Figure 2b. Fast-Fourier transform patterns are included as insets: (a) hcp-Re 1�101��orientation; (b) fcc-Niat [001]orientation;and(c)hcp-Re 2�423��

orientation.Thephases in(a) and (b) share their (001)fcc and ð01�11Þhcp stacking planes with inter-planar spacings of1.917 and 2.107nm, respectively, in accordance with the SAED pattern, Figure 2c.


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results of Cohen-Sagiv et al.[9] of a 10�10Þ�and 01�11Þ�

preferredorientation, perpendicular to the surface of a Cu substrate, ofhcp-Ni and hcp-Ir0.4Re0.6 phases, based on X-ray diffraction(XRD) data.


Figure 4 displays a STEM-EDS line scan, white arrow inFigure 2a, through the dark and bright layers. Table 1compares the atomic concentrations of elements as obtainedby SEM-EDS analyses of the top surface, STEM-EDS analysesof distinct layers, and LEAP tomography analyses within thecolumnar grains and at the intercolumnar GB. As is evidentfrom the STEM-EDS analyses, the concentration of Re in thebright layers is greater than in the dark layers (58.6 at% vs.49.5 at%, respectively). These local compositions match theZ-contrast observations in the STEM-HAADF image,Figure 2a. Additionally, given the sampling interactionvolume in SEM-EDS and the significantly higher volumefraction of the Re-rich layers, the deviation of the chemicalcomposition of the bright layers as measured by STEM-EDSfrom the nominal composition determined by SEM-EDS issignificantly smaller than that of the dark layers. Thecomposition of the dark layers was found to be similar tothat at the intercolumnar grain boundaries. Additionally,except of traces of Ga and Pt from the ion-beam milling andprotective coating procedures in FIB, respectively, no otherimpurity atoms were found in the coating.

3.4. Three-Dimensional (3D) Chemical Mapping byLEAP Tomography

Figure 1a demonstrates that the surface crevices atintercolony boundaries may serve as preferred sites for crackpropagation along intercolumnar GBs, due to the two types ofboundaries trying to remain continuous. This implies thatthe chemical composition of the grain boundaries may bedifferent from that in the colonies. Similarly, the chemicalcomposition at the intercolumnar GBs may be differentfrom that within the grains. Thus, the elemental distribu-tions at the intercolumnar GBs were determined by APT.Figure 5a displays the STEM-HAADF image of an APTnanotip that was prepared such that it would contain anintercolumnar GB. Figure 5b is a 3D LEAP tomographicreconstruction of field-evaporated atoms displaying nickel,



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rhenium, oxygen, hydrogen, and rhenium oxide atoms (a full3D movie is available in the Supporting Information S2). TheGB region is highlighted using a green isoconcentrationsurface, corresponding to 52 at% Ni. The chemical composi-tion at the intercolumnar GB can be determined by drawing aline profile between two adjacent grains, Figure 5c. Theconcentration profile exhibits an asymmetry because of therough interface on the right-hand side of a columnar GB.Based on this figure, the intercolumnar GB contains 52.4 at%Ni, which is greater than within the grains. This compositionis similar to the one of the dark layers as analyzed by STEM-EDS. Themeasurement of hydrogen atoms and light elementsis one of the many strengths of APT.[43–46] The importantfeature of an APT is that it is a spectrometric technique using amicrochannel plate (MCP) for detecting ions and measuring

Fig. 5. Atomic-resolution chemical analyses by LEAP tomography: (a) STEM-HAADFimage of an APT microtip, which contains an inter-columnar grain-boundary; (b) APT3D reconstruction based on field-evaporated ions (�3.2 million ions) from the APTnanotip. Nickel (small green dots), rhenium (small red dots), oxygen (cyan spheres),hydrogen (blue spheres), and rhenium oxide (dark red spheres) are included. The grain-boundary region is highlighted in green using 52 at% Ni isoconcentration surfaces. (c)Elemental concentration profiles along a line that crosses the inter-columnar GB. Notethe different ordinate scales for Re and Ni contrasted with O and H.


their mass-to-charge state ratios (m/n), essentially one ion at atime.[21–23] Once the energy of an ion is greater than a fewkilovolts, the detection efficiency of an MCP is essentially thesame for all elements in the periodic table. Due to the highmass resolving power of the APT for all elements in theperiodic table, hydrogen and oxygen are readily distin-guished from the heavier refractory metals. A quantitativeAPT study of hydrogen in ultrahigh purity niobium thatcontains b-niobium-hydride (b-NbH) precipitates has beenreported.[47,48] The hydrogen concentrations (H/Nb atomicratios in the range of 0.7–1.1) were found to be in reasonableagreement with the expected hydrogen concentrations. Whilethe concentrations of oxygen and hydrogen at the inter-columnar GBs are greater than at the bright layers within thegrains, they are still significantly smaller than those typical ofoxides and hydrides. Thus, the coating material may beregarded as an alloy. Nevertheless, the concentration ofhydrogen in the Re-rich layer (0.4 at%) and at theintercolumnar GB (1.2 at%) is sufficiently high to indicate apossible effect of hydrogen on the texture and grain size of theRe–Ni electrodeposits and the intercolumar GB cracking thatis most likely associated with hydrogen absorption, hydrideformation, or hydrogen embrittlement.

4. Discussion

4.1. Intrinsic Stresses Developed in the Banded Structure ofthe Re–Ni Electrodeposit

The formation of alternating layers of Re-rich and Ni-richtwo phases requires an evaluation of the Re–Ni binary phasediagram, Figure 6.[49–51] This is a peritectic phase diagramwith an extensive region of mutual solubility and nointermetallic compounds. Electrodeposition, by definition,is a non-equilibrium process as it occurs under a cathodicoverpotential. Consequently, the phases formed are often notin thermodynamic equilibrium, and may be either undersat-urated or supersaturated compared to their counterparts inthe equilibrium phase diagram.[52] Even in electroplating,

Fig. 6. The Re–Ni binary phase diagram.[51]



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binary alloys with a peritectic (or eutectic) phase diagram
often crystallize into two types of crystals, with the highermelting point phase being fine-grained.[53] Comparing thechemical compositions listed in Table 1 to the phase diagram,it is clear that the Ni-rich layer is supersaturated in Re,whereas the Re-rich layer is supersaturated in Ni.

Combining the APT data with STEM-HAADF measure-ments, we estimated the intrinsic stresses developed inthe banded structure of the Re–Ni electrodeposit. Thelattice parameters for hcp-Re and fcc-Re are a¼ 0.2761 nm,c¼ 0.4456 nm,[42] and a¼ 0.393� 0.004 nm,[54] respectively.While hcp is the only stable phase of pure Re, a metastablefcc-Re phase has occasionally been observed, either experi-mentally (e.g., by high undercooling[54] or epitaxial growthon a NaCl substrate[55]) or in atomistic calculations.[56] Thelattice parameters of fcc-Ni and metastable hcp-Ni area¼ 0.3524 nm,[49] and a¼ 0.2622 nm, c¼ 0.4321 nm,[57] respec-tively. Cohen-Sagiv et al.[9] reported the formation of the hcp-Ni phase in Re–Ir–Ni electrodeposits based on XRD data. Onthe basis of the results reported herein, the formation of themetastable hcp-Ni phase may be governed by the high strainsassociated with the hetero-epitaxial growth of Ni-rich layerson hcp Re-rich layers. Utilizing the values of the latticeparameters and the chemical composition of the layers,Table 1, the predicted lattice parameters based onmixing of ReandNi are a¼ 0.2712 nm, c¼ 0.4403 nm for the Re-rich (bright)phase, and a¼ 0.2683 nm, c¼ 0.4378 nm for the Ni-rich (dark)phase. The TEM data in the current research and the XRDdata reported by Cohen-Sagiv et al.[9] demonstrate thathetero-epitaxial growth along the normal to the substrate isparallel to the 01�11��

and 10�10��crystallographic directions,

respectively. The strain developed in the Ni-rich phaseresulting from the hetero-epitaxial growth is e¼ 0.0057.An effective shear modulus of 105.6GPa was calculatedbased on the shear moduli of Ni and Re (76 and 155GPa,respectively). Thus, the residual shear stresses developed inthe Ni-rich phase are �602MPa. This large value can explainsome of our microscopic observations, e.g., the formation ofrough interfaces between the Ni-rich and Re-rich layers,cracking, etc.

4.2. The Essence of the Formation of a CombinedColumnar-Multilayer Structure

While columnar structures and banded structures haveoften been observed in electrodeposits, a combination of bothin the same coating is rare; specifically, in coatings processedfrom a single bath by dc electroplating. Thus, the final part ofthis article focuses on reviewing the relevant literature andsuggesting a deposition mechanism for the Re–Ni system.

The initial electrodeposited layer inherits the structure ororientation of the substrate when the misfit strain energybetween this layer and the substrate is less than 12%.After thisepitaxial layer achieves a critical thickness (1mm or less), finecolumnar grains extending along the growth direction mayform and the deposited crystallographic texture then becomesindependent of the substrate structure. These microstructures


are strongly dependent on the different deposition conditions,such as the plating bath composition and temperature,applied current density, orientation and surface morphologyof the substrate, etc.[58] Several theories have been proposed toexplain the development of microstructures in electrodepo-sits. The microstructure often appears to be the outcome of acombination of different mechanisms. The electrocrystalliza-tion process is influenced by several components of the overalloverpotential, namely the activation overpotential, theconcentration overpotential, the resistance overpotential,and the crystallization overpotential.[37,59,60]

The structures of electrodeposits are mainly classified as:i) fine-grained (mostly equiaxed); ii) columnar; iii) fibrous; oriv) banded structure.[39,61] Columnar structures are character-istic of deposits from additive-free solutions (especiallyacidic), which are formed at a low current density (i.e., lowdeposition rate) or elevated temperatures. They are typicallycharacterized by lower tensile strengths and microhardnessvalues than other structures, but are generally more ductileand have high purity and density, and a small electricalresistivity. In contrast, banded structures are characteristicof brighteners-containing solutions, and in processes thatinvolve either multiple baths or periodic variations of somedeposition parameter (e.g., pulse plating). Spontaneousformation of banded structures have, however, also beenobserved in the electrodeposition of some pure metals (e.g.,Ni) and some alloys (e.g., Ni–P and Ni–Co).[62] The layersoften have a sub-micron thickness, and are thought tobe compositional or structural modulations that originatefrom small-scale periodic fluctuations in the metal oradditive concentrations, pH, agitation, or polarization atthe cathode’s surface.[37,62] Banded structures generallypossess higher tensile strengths and microhardness values,but also higher internal stresses and lower ductility than otherstructures. Additionally, compositionally modulated multi-layers have been reported to exhibit remarkable wearresistance, magnetic, electronic, superconducting, and opticalproperties.[63]

In Figures 1c and 2a, corrugated layers are evident. Rafajaet al.[64] reported the effect of current density on the structureof Co–Cu/Cu multilayers electrodeposited from a single bathunder the potentiostatic/galvanostatic mixed mode. At asufficiently low current density (20.7mAcm�2), both colum-nar grains and multilayers were formed. Highly correlatedinterfacial corrugations were observed and related to theinterplay between local lattice strains and segregation ofdifferent elements. Using a continuum mechanics approach itwas shown that under thermodynamic equilibrium condi-tions there is a critical film thickness below which misfit(edge) dislocations do not form; as the lattice parametermismatch decreases, the critical film thickness for misfitdislocation formation increases.[65] Spaepen reviewed theeffects of interfaces on the intrinsic stresses in polycrystallinevapor-deposited thin films, including multilayered films.[66]

Explanations were given for phenomena such as layercurvature and rough interfaces observed in this study. It


Fig. 7. The variation of potential with time during electrodeposition under thegalvanostatic condition (i¼ 50mA cm�2). (a) High potential spikes are evident duringthe first 138 s: see the full time and potential scales. (b–d) Plots for three different timeregimes within (a), which reveal periodic fluctuations in potential (polarization) aswell as different potential profiles corresponding to different microstructures observedby TEM.


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was shown, e.g., that the non-planarity of an interface candecrease the value of the interface stress by as much as 60%(for a sinusoidal variation of roughness and if the amplitudeof the variation is equal to the wavelength). Additionally,Bruner and Kern reviewed the effects of strain on hetero-epitaxial metal growth.[67]

The effects of the adsorption of hydrogen atoms on thetexture of the coating have been suggested for columnar[68]

and lamellar growth[69] during electrodeposition of Ni. Freegrowth can be related to a substrate surface in the absence ofadsorbed hydrogen atoms and surface-active substances. Forsingle crystals, different lattice planes have different freegrowth metallic overpotentials; the higher the reticulardensity (number of atoms per unit area on the face), thelower is the growth rate in a direction normal to the face.Hydrogen atom intermediates are adsorbed more easily onlattice planes, which during free growth have faster growthvelocities. Thus, hydrogen adsorption results in a greater freeenergy of activation for Ni deposition, i.e., an increase in themetallic overpotential above the free growth value. Thetexture is affected by the average surface coverage ofhydrogen atoms (uav); the higher the value of uav, the higheris the hydrogen content of the electrodeposit. Different valuesof uav result in different textures.

4.3. Electrochemical Mechanism for Re–Ni MultilayerCoating

Finally, we now describe the microstructural evolution inthe Re–Ni coating and discuss the electrochemical mechanismfor its formation, based on our TEM and APT results. Severalexperimental conditions of deposition were changed com-pared to our prior studies,[5–7,9] including lowering the pH to2.3–2.4 (instead of either 5.0 or 8.0), using a stainless steelsubstrate (instead of pure Cu), and adding an activation stepin the Pd-citrate buffer solution prior to deposition. Whilethese changes were aimed at either improving the quality ofthe coating or making it easier to manipulate the specimens inthe dual-beam FIB microscope, they may affect both theelectrochemical reactions involved in the deposition and themicrostructure of the coating. Stainless steel substrates aregenerally regarded as being covered with a passive amor-phous oxide layer, which has no epitaxial effect on the initialdeposit.[53] Moreover, the Pd-activation step provides manynucleation sites for the initial deposit. Thus, an initiallydeposited layer, consisting of fine nanograins and having athickness of less than 500 nm, is first formed on the surface ofthe substrate. In the early stage of deposition, agitation in thevicinity of the cathode’s surface due to hydrogen evolution isstill relatively small, which enables the formation of equiaxednanograins. In the second stage of deposition, columnargrowth occurs from the acidic, additive-free plating bath,indicating that the deposition rate is fairly small. Simulta-neously, compositionally modulated multilayers are formedwithin each columnar grain due to the peritectic character ofthe Re–Ni system and small-scale periodic fluctuations in theconcentrations of nickel and perrhenate ions and the pH in


the vicinity of the cathode surface, and accordingly in thereversible electrode potential and the overpotential. Figure 7presents experimental proof of these periodic fluctuations,which displays changes in the electrode potential versus timeduring deposition, which was performed under the galvano-static condition. High-potential spikes are evident within thefirst 138 s of deposition, Figure 7a. Based on the proportions oftime and layer thickness, we hypothesize that these largefluctuations in potential are associated with the formation ofthe initial layer. Subsequently, the potential appears to be in asteady-state, at low magnification. Figure 7b–d demonstratesclearly fluctuations in potential that are not periodic duringthe entire deposition process. The potential profile for150–800 s, Figure 7b, is different from that for 800–2600 s,Figure 7c. The latter exhibits potential transients in the form ofpotential steps, with three plateaus between �1.37 and �1.46VAg/AgCl, which presumably include the potential profileduring deposition of the center layer in which the columnargrains are smaller. During the last 800 s of deposition,Figure 7d, there are several significant potential increases,but within each mean plateau value the profile is less noisy. Itis during this time regime (2800–3600 s) that the top nodular(colony) structure was formed.

High internal tensile stresses developed in the coating dueto co-deposition of hydrogen along with the metallicelements, hetero-epitaxial growth and, possibly, decomposi-tion of a hydride phase. These stresses result in formation ofmisfit dislocations, cracking, corrugations of layers, anddiscontinuous layers when crossing an intercolumnar GB inthe lateral direction, and possibly formation of metastablephases, such as fcc-Re and hcp-Ni.


Fig. 8. Schematic representation of the mechanism of electrodeposition of the Re–Nicoating. Freshly deposited Ni0 (a) catalytically reduces the perrhenate ion ReO�

4 to the5-valent oxidation state in the rhenate ion ReO�

3 , or directly to Re0. The nickel ion acts

as a catalyst (b). Possible routes for deposition of Ni are suggested elsewhere.[10]


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At pH 2.3–2.4, where electrodeposition was performed, the
relative abundances of the two major citrate-containingspecies are 81.3% H3Cit and 18.5% H2Cit

–. The former isnot an effective complexing agent.[2] The actual pHwithin thediffusion layer, i.e., in the solution adjacent to the cathode, wasprobably fluctuating during deposition, which could be thecause of multilayer formation. Furthermore, APT revealedhigher concentrations of hydrogen and oxygen in the Ni-richphase, which contains intercolumnar grain boundaries,compared to the Re-rich layers. The high Ni concentrationsat the boundaries between grains of the Re-rich phase indicatethat it behaved similar to foreign substances, which are sweptahead of a growth front during electrodeposition. Moreover,for Ni electrodeposition, it has been suggested that adsorptionof a nickel hydroxide intermediate and the release of only partof the hydroxide ions into the solution results in the occlusionof oxygen and hydrogen (as part of nickel hydroxideprecipitates) in the deposit. The precipitation of nickelhydroxide was enabled due to a significant increase inthe pH of the solution adjacent to the cathode.[70]

The formation of a banded Re–Ni structure may indicatealternate depletion and enrichment of nickel ions in thediffusion layer. When a Re-rich layer forms, more hydrogen isbeing formed caused by the catalytic effect of Re on thehydrogen evolution reaction. Consequently, the pH isincreased locally, thus favoring the formation of a nickelhydroxide intermediate at the surface of the cathode.Hydrogen evolution can also result in more nickel ions inthe diffusion layer due to increased convection. When a Ni-rich layer forms, less hydrogen is evolved, the local pHdecreases, and a new cycle begins. The bands of the Re-richphase are thicker than the bands of the Ni-Rich phase. Thisimplies that the spacing between Re-rich layers is smaller thanthe spacing between Ni-rich layers. Thus, it is argued that thedepletion of nickel ions in the diffusion layer is less significantthan the depletion in perrhenate ions. This argument furthersupports the mechanism Naor et al. have suggested for theco-deposition of rhenium with iron-group metals.[5–7] Theessence of this mechanism is illustrated in Figure 8.

It is proposed that reduction of the divalent Ni2þ ion,formingmetallic Ni0 on the surface of the cathode, is followedby chemical reduction of the heptavalent rhenium in the stableperrhenate anion into pentavalent rhenium in the less stablerhenate anion. The latter can be reduced either electrochemi-cally or chemically, in several steps, to metallic Re0. Thechemical reduction of heptavalent rhenium is accompanied byoxidation of Ni0 to Ni2þ, and the release of the latter back intosolution. Thus, nickel may be regarded as a catalyst, or areducing agent that is formed in situ. One of the remainingopen questions concerns how Ni is incorporated in thecoating; is it deposited in parallel to Re, forming a Re–Ni alloy,or does it simply act as a catalyst that is partially beingoccluded in the deposit because the rate of its re-oxidation istoo slow compared to the rate of Re deposition.[6,7] Based onthe correlative APT and STEM results of our study, it appearsthat Ni is actually co-deposited as an element of a binary alloy.


4. Conclusions

This study involves atomic-scale structural and atomicpart-per-million level chemical characterization of a Re–Nibinary alloy coating. Toward this goal, aberration-correctedscanning transmission electron microscopy (STEM) andatom-probe tomography (APT) analyses were utilized forRe-based coatings. A unique combination of a columnar andmultilayer structure is observed, for the first time, for anelectrodeposit produced by dc plating in a single bath. Themultilayers exhibit a hetero-epitaxial growth behavior,which results in the development of very high residualshear stresses (�600MPa). These high stresses are most likelyresponsible for several phenomena that are reportedherein, including formation of rough interfaces between theNi-rich and Re-rich layers, cracking in the outer layer, etc.AlternatingRe-andthinnerNi-rich layerssupportamechanisminwhichnickel acts asa reducingagent, allowingthedepositionof rhenium. These observations shed new light on the role ofiron-group metals in catalyzing the electrodeposition ofrhenium and increasing significantly the Faradaic efficiency.Simultaneously, this study also calls for the development ofalternative routes for high-quality electrodeposition of essen-tially pure rhenium coatings; thereby, modifying the structureof the coating material, reducing residual stresses and



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eliminating micro-cracking, which will be of paramounttechnological importance given the great interest in rhenium-based coatings for high-temperature, catalytic and otherapplications.

Article first published online: March 29, 2016Manuscript Revised: March 1, 2016

Manuscript Received: November 16, 2015

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