Surface & Coatings Technology 205 (2011) 3651–3657

Contents lists available at ScienceDirect

Surface & Coatings Technology

j ourna l homepage: www.e lsev ie r.com/ locate /sur fcoat

Influence of lower current densities on the residual stress and structure of thicknickel electrodeposits

Siddhartha Pathak ⁎,1, Marion Guinard 1, Martine G.C. Vernooij, Barbara Cousin 2, Zhao Wang 3,Johann Michler, Laetitia PhilippeEmpa, Swiss Federal Laboratories for Materials Science and Technology, Feuerwerkerstrasse 39, CH-3602 Thun, Switzerland

⁎ Corresponding author. Present address: MaterialsTechnology (Caltech), 1200 E. California Blvd., MC 309-USA. Tel.: +1 626 395 8165; fax: +1 626 395 8868.

E-mail addresses: sp324@drexel.edu (S. Pathak), ma(M. Guinard), vernooij@alumni.ethz.ch (M.G.C. Vernooi(B. Cousin), wzzhao@yahoo.fr (Z. Wang), johann.michlelaetitia.philippe@empa.ch (L. Philippe).

1 Both authors contributed equally to this work.2 Present address: LISI Automotive Former, Rue Ferna

sur Cher, France.3 Present address: LITEN/LCH, CEA-Grenoble, 17 Rue

France.

0257-8972/$ – see front matter © 2011 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2011.01.012

a b s t r a c t

a r t i c l e i n f o

Article history:Received 6 August 2010Accepted in revised form 6 January 2011Available online 14 January 2011

Keywords:NickelElectrodepositionSulfamate bathTextureAnisotropyBack-scattered electron diffraction (EBSD)Indentation

Thick Ni electrodeposits, with thicknesses ranging from 18 to 90 μm, have been produced from sulfamatebaths using very low current densities (1.5–5 mA cm−2). Increasing grain sizes, with a correspondingdecrease in the hardness values, and more columnar grain morphologies were observed with increasingcurrent densities in the deposits. The anisotropy in the mechanical properties of the deposits were found tocorrelate strongly with the deposit texture, where the change from a b101N dominated orientation at1.5 mA cm−2 to a b001N orientation at the higher (5 mA cm−2) current density led to a correspondingdecrease in the indentation modulus values for the respective Ni films. The residual stress values measured inthe deposits (75–136 MPa) were found to be comparable to literature values of deposits plated at highercurrent densities, which rules out any potential advantage from the use of such low current densities forproducing lower residual stresses.

Science, California Institute of81, Pasadena, CA 91125-8100,

rion.guinard@empa.chj), cbab.s@hotmail.frr@empa.ch (J. Michler),

nd Léger, 18400 Saint Florent

des Martyrs, Grenoble 38000,

l rights reserved.

© 2011 Elsevier B.V. All rights reserved.

1. Introduction

Electrodeposited nickel, particularly using sulfamate electrolytes,are widely popular in the electroplating andmicrofabrication industrydue to their high deposition rates, greater film thickness, goodmechanical and corrosion resistant properties, and low residual stress[1–5]. Recent reports have also suggested the use of thick (~tens orhundreds of microns) nickel films as matrix for the production ofnickel–diamond composite coatings for high precision wear-resistantcomplex shaped grinding or cutting tools [6–8]. However, like mostelectrodeposited metals, nickel suffers from some common problems,such as anisotropic material properties, poor repeatability, and highsensitivity to electroplating conditions [3,4]. For example, an increasein current density is often used to reduce the yield time and increasethe output of these electrodeposits for greater film thickness [1,2,9],and thus most of the reported investigations on Ni electrodeposits

have been performed on films made using high current densities inthe range of 5–40 mA cm−2 [2,3,9] or higher [1]. However a highercurrent density can be detrimental for the mechanical properties ofthe thick deposits due to structural changes occurring in the film,deformation of the tools [6], as well as residual stresses developedduring the processing operations [5,10]. On the other hand, lowerdeposition rates (current density) are expected to produce depositswith improved mechanical properties along with lower tensilestresses for films thicker than 1 μm [3,11]. As such lower currentdensities are often used in the electroplating industry in order toproduce stress-free deposits, although such a practice is detrimentalto the production speed. The focus of the current paper is to verify therelationship between residual stress and current density using Nielectrodeposits over a range of thicknesses (18–90 μm) made usingunusually low current densities (1.5–5 mA cm−2).

Both grain geometry and grain size, as well as the crystallographictexture of electrodeposited Ni microstructures, are known to dependstrongly on the plating conditions [2]. Thick Ni electrodeposits platedat lower current densities (3–15 mA cm−2) have been reported tohave a characteristic columnar microstructure where the grains oftenextend through the thickness of the deposit, with the grain sizeincreasing with increasing current densities [9,12]. The surfacemorphologies of the deposits have also been shown to change froma ‘cauliflower’ to a ‘truncated pyramidal’ structure with increase(N10 mA cm−2) in current densities [13,14]. The crystallographictexture for thick (~microns) Ni electrodeposits is known to beindependent of the substrate material; instead it is found to dependstrongly on the current density, pH, and presence of organic additives

3652 S. Pathak et al. / Surface & Coatings Technology 205 (2011) 3651–3657

in the electroplating solution. Two general textures have beenreported for electrodeposited Ni: a strong b100N texture caused bya regime of ‘free growth’ at higher current densities (≥10 mA cm−2),and a b110N texture at lower current densities (≤5 mA cm−2) whereNi growth is inhibited due to impurities and/or hydrogen adsorptionon the plating surface [3,9]. Indentation hardness values for electro-deposited Ni are also generally found to decrease with increasingcurrent densities [2,3].

The highly textured-oriented grain morphology in electrodepositedNi suggests a significant anisotropy in the mechanical response of thismaterial, which needs to be considered in detail [15–17]. Indeed, asignificant range of mechanical properties have been reported inliterature for electrodeposited Ni [2,3,12,16], which could have beencaused inpart by simply averaging the response fromdifferent locationsof the polycrystalline sample without accounting for their underlyingorientations. It is nowpossible tomeasure the local lattice orientation inpolycrystalline samples using a technique called Orientation ImagingMicroscopy (OIM) [18,19], which is based on automated indexing ofback-scattered electron diffraction (EBSD) patterns (obtained using ascanning electronmicroscope) and has a spatial resolution of less than amicron. Therefore, a judicious coupling the structure informationobtained from OIM with the mechanical data obtained at the samelength scale (using nanoindentation) could provide interesting insightsabout the anisotropicmechanical response of the local Ni film structure.

The present work is an experimental study on the residual stresses,and the resulting structural and mechanical changes, occurring innickel electrodeposits produced using very low current densities. Weutilize the combined capabilities of profilometry, nanoindentationand EBSD to study the anisotropic microstructure–property relation-ships in electrodeposited nickel in light of the above observations.

2. Experimental techniques

2.1. Materials and processing

Microcrystalline nickel electrodeposits were produced by directcurrent (DC) plating in a nickel sulfamate electrolyte with 1.5 g/L ofnickel chloride, 1.5 mL/L of wetting agent and 35 g/L of boric acid.Three current densities were chosen in order to study their effect onresidual stresses developed in electrodeposited Ni: 1.5, 2.5 and5.0 mA cm−2. Note that these current densities are lower than thetypical range (5–40 mA cm−2) generally studied in literature [2,3,9].The electrodeposition was conducted at 30±1 °C, without stirringand at a pH of 3.5. Nickel balls in a plastic basket were used as anode.The substrate was an aluminium alloy 2017 (1 mm thick plates with18 cm2 area). Before electrodeposition, the substrate was cleaned,deoxidised, activated by phosphatizing and finally prepared by anickel electroless depositionmechanism. The deposition time for eachcurrent density was varied to keep the total charge constant accordingto Faraday's law [20], thus enabling the same nominal thickness foreach set of measurements. The average deposit thicknesses, asmeasured by profilometry after electrodeposition, were found to be18, 69 and 90 μm. The current efficiencies were between 80 and 90%.The surface topographies of the deposits were examined using a ZeissDSM 962 scanning electron microscope (SEM), operated at 20 kVaccelerating voltage. Element profiling by Glow Discharge OpticalEmission Spectroscopy (GDOES) showed an average of 95 at.% Ni inthe deposited films, with oxygenmaking upmajority of the remaining5 at.%.

A detailed study of themicrostructure–property correlations in theNi deposits was carried out on the largest nominal thicknesses of thedeposits (90 μm), chosen for the ease of their surface preparation. Forthese tests, the samples were cut into ten pieces of 1.8 cm2 with acircular saw, and the center pieces were chosen for further analysis.After the cutting, the samples were cleaned ultrasonically by a three-step cleaning process using acetone, ethanol and water. They were

polished to a high surface finish (0.25 μm diamond suspension) andsubsequently subjected to lapping for 5 minutes using an alkalinecolloidal silica suspension. Finally the samples were polished with Arions at 4 kV and 2.5 mA for 1 hour (angle of incidence 5°) using aBaltec RES 010 device.

The EBSD patterns were obtained using 20 kV acceleration voltageat 15 mmworking distance and 69 μA beam current with a 70° sampletilt. Orientation maps were collected using a hexagonal grid of pointswith a step size of 150 nm and typical collection speed of 18.5 patternsper second. The patterns were collected, indexed and analyzed withthe commercial EDAX/TSL software OIM™ 5.2. Since this technique ofautomatic pattern indexing probes only the very top layer of thesample, orientationmaps were collected on the deposits' top surfaces,as well as along their cross-sections, in order to evaluate theirstructure both normal and parallel to the growth direction, respec-tively. Grain sizes along both these directions were calculated usingthe orientation maps [21].

2.2. Mechanical measurements

Nanoindentation studies were carried out on the top surface of thedeposits in order to correlate the structure information obtained fromOIM-EBSD with the mechanical properties of the deposits. All testswere performed using a nanoindenter (MTS XP® System) with aBerkovich diamond indenter tip under load control to peak displace-ments of 2000 nm at 50 nm/s loading–unloading rate with a 10 s holdbefore unloading in all samples. At least 10 tests were carried out foreach sample. Since this indentation depth is significantly lower thanthe thickness of the deposits, the mechanical data obtained fromnanoindentation can be taken to be representative of the properties ofthe deposits without any substrate influence. The modulus andhardness of the deposits were obtained from the nanoindentationload–displacement curves by analyzing the initial (50%) unloadingsegment using the Oliver and Pharr method [22]. The projectedcontact area for indentation was obtained by calibration on fusedsilica, in order to account for deviations from non-ideal indentergeometry.

Note that the Oliver and Pharr method is largely based on Hertz'stheory [23,24], and as such addresses only the elastic indentation of anisotropic sample [25]. However, in practice, the elastic response of thesample at the typical length scale of nanoindentation is ofteninherently anisotropic [26], especially in polycrystalline metals suchas Ni where the indents are comparable to or smaller than the typicalgrain size. In order to analyze the elastic indentation of anisotropicsamples, the treatment suggested by Vlassak and Nix [27–30] wasused in this work. These authors suggest that a modified Hertz theorycan be used for elastic indentation of cubic crystals, provided ananisotropy parameter, β, is appropriately introduced into thedefinition of the effective indentation modulus as given below:

1Eeff

=1β

1−ν2sample

Esample

!+

1−ν2indenter

Eindenter

!ð1Þ

where Eeff denotes the effective Young's modulus of the indenter andthe specimen system, Esample and νsample(=0.31 [31–33]) denote theeffective values of Young's modulus and Poisson's ratio, respectively,for a randomly textured polycrystalline aggregate of crystals with thesame elastic properties as the single-crystal being studied [28,29], andνindenter (=0.07)and Eindenter (=1141 GPa) denote the Poisson's ratioand theYoung'smodulus for the indentermaterial (diamond). For cubiccrystals, the value ofβdepends strongly on the crystal lattice orientationand the degree of cubic elastic anisotropy. The elastic anisotropy (A) of acubic crystal is usually defined as A=2C44/(C11−C12), where C11, C12,and C44 denote the cubic elastic constants used to define the crystalelastic stiffness in its own reference frame. Based on the values reported

Fig. 1. SE images showing the surface morphologies of the 69 µm thick nickel elec-trodeposits produced at three current densities: (a) 1.5 mA cm−2, (b), 2.5 mA cm−2,and (c) 5 mA cm−2.

3653S. Pathak et al. / Surface & Coatings Technology 205 (2011) 3651–3657

by Vlassak and Nix [28,29], β should range from 0.9 to 1.05 (a rangeof ~15%) for Ni crystals (for which A=2.45 [31,32]) of differentorientations. Note that this range is around 10% if the orientations varyonly from (100) to (101). In this work, the modulus values for ourelectro-deposit samples (Esample) were corrected for anisotropy usingthe values of β as reported in [28,29].

During electrodeposition metal thin films typically grow undernon-equilibrium conditions, which often result in internal (orresidual) stresses even in the absence of any external force ortemperature gradient [34]. The Stoney equation [35], or one of itsvariants, is typically used to experimentally determine such stressespresent in a thin film. However, the inherent assumptions of theStoney formula – such as a negligible thickness of the film ascompared to the substrate and the lateral dimensions – are notapplicable for the thick films developed in this work. Here we haveused a modified version of the Stoney equation, shown by Clyne[36,37], to relate the measured curvature to the stress in our thickfilms, as shown below:

σR = Mf κE2f t

4f + 4Ef Est

3f ts + 6Ef Est

2f t

2s + 4Ef Estf t

3s + E2s t

4s

6EsEf ts + tf� �

tstfð2Þ

Here, Mf=Ef/(1−υf) is the effective biaxial modulus of the film,κ=1/R−1/R0 is the curvature, R0 and R are respectively the samplecurvatures before and after electrodeposition, υ, E and t are thePoisson's ratio, Young's modulus and thickness, and the subscripts sand f refer to the substrate (aluminium alloy 2017) and the film,respectively. The Young's moduli of the Ni films and the Al alloysubstrate were measured by indentation (Esample) using Eq. (1) asdescribed above. The indentation modulus for the Al alloy substratewas measured to be 79 GPa, which is close to the reported modulusvalue for this alloy (73.2 GPa [38]). A value of 0.31 was used for thePoisson's ratio (νs) of the film [31–33]. The sample curvatures weredetermined using a white light surface profilometer ALTISURF 500with a 3 cm distance between sample and light spot, both before andafter electrodeposition. The substrate thicknesswas 1000±1 μm in allcases.

3. Results

3.1. Grain morphology

SEM images of the 69 μm thick deposit surfaces (Fig. 1) show thatan increase in the current density from 1.5 to 5 mA cm−2 induces achange in the grain morphology of the Ni electrodeposits. Nideposited at the lowest current density of 1.5 mA cm−2 (Fig. 1a)appears to have a “cauliflower” [14] or “colony” [13] structure, withdeep crevices outlining groups of smaller substructures. Increasingthe current density causes the morphology to slowly change to a“truncated pyramidal type” structure [13,14] (Fig. 1b and c). The sizesof the surface structures are also seen to increase with increasingcurrent density.

A change in the current density is also seen to greatly influence thecross-sectional grain morphology, as seen from Fig. 2 for the 90 μmthick deposits. The EBSD maps of the surface and the cross-section inthis figure show that the structure of the deposits evolves into a morecolumnar structure at higher current densities. At the highest currentdensity of 5 mA cm−2

, some grains can be seen to extend almostthrough the thickness of the deposit. There is also an increase in thegrain size as the current density increases from 1.5 to 5 mA cm−2.Both these effects are summarized by the grain size measurementsnormal and parallel to the plating/growth direction in Table 1.

3.2. Texture

The structure–property correlations from the OIM-EBSD andnanoindentation studies are summarized in Fig. 2 and Table 1 forthe 90 μm thick deposits. In Fig. 2, the grain orientations are color-coded to reflect their positions in the inverse pole figure map. Thus,the grains that have a {001} crystallographic plane parallel to thesample surface are colored red. Note that since Ni is cubic, this alsomeans that the grains colored red have a b001N crystallographic

Fig. 2. (Color online) Texture and grain morphology of the (a) sample surface and (b) sample cross-section for the 90 μm thick Ni deposits at three different current densities. Thegrain orientations in the OIM-EBSDmaps have been color-coded to reflect their positions in the inverse pole figure (IPF)map. The orientation density IPF contour maps in the bottomrepresent the densities of various orientations as compared to a standard sample of random orientation (the orientation density for a random orientation is equal to 1).

3654 S. Pathak et al. / Surface & Coatings Technology 205 (2011) 3651–3657

direction parallel to the sample normal direction (ND; also theindentation direction). Similarly, grains with {101} and {111} planesparallel to the surface are colored green and blue respectively. Thecrystallographic textures within the cross sections of the deposits are

Table 1Summary of themicrostructure–property correlations for the 90 μm thick Ni deposits atthree different current densities. The residual stresses were calculated using Eq. (2),while the grain sizes were calculated from the OIM-EBSD maps (Fig.2). Since the grainsize distributions are skewed towards the smaller grains in all of the samples, thestandard deviations in grain size measurements are relatively large. The hardness andeffective modulus values (Eeff) were calculated from nanoindentation [22]. The valuesfor sample modulus (Esample) were corrected for the anisotropy parameter β usingEq. (1).

Current Density 1.5 mA cm−2 2.5 mA cm−2 5.0 mA cm−2

Residual stress 74.5 MPa 83.2 MPa 94.1 MPaHardness 3.46 GPa 2.74 GPa 2.49 GPaGrain size (surface) 0.8±0.4 μm 1.1±0.8 μm 1.8±1.5 μmGrain size (parallel togrowth direction)

1.5±1.4 μm 2.0±2.3 μm 3.9±6.1 μm

Effective modulus (Eeff) 186.3±15 GPa 179.5±9 GPa 176.2±14 GPaSample modulus (Esample) 214.5±14 GPa 213.9±9 GPa 213.9±14 GPa

shown in inverse pole figure (IPF) density distributions, obtained by aGaussian convolution with a half width of 5° within the harmoniccalculus (bin size 5°) [19]. In the IPF density distribution maps, the IPFfor a random standard sample is designated as ‘×1’ of all points. Thus,all points marked ‘×4’ indicate an IPF density of four times random.

It can be seen from Fig. 2 that the current density has a markedeffect on both the grain orientation as well as the grain morphology.At the lowest current density (1.5 mA cm−2) the texture is dominatedby the {101} orientation along the plating direction, which slowlychanges to favor the {100} planes at the higher (5 mA cm−2) currentdensity. The {101} texture is still dominant at the intermediatecurrent density level (2.5 mA cm−2), but small regions with {100}grains can also be seen here.

3.3. Nanoindentation modulus and hardness

Current density also has a significant influence on the hardnessand modulus values of the deposits, as measured by nanoindentationusing the Oliver and Pharr method [22]. As seen from Table 1 for the90 μm thick deposits, the hardness of the deposits slowly decreasesfrom 3.46 to 2.49 GPa with increasing current density. A similar trendis seen for the modulus measurements, where the average effective

3655S. Pathak et al. / Surface & Coatings Technology 205 (2011) 3651–3657

modulus (Eeff) is seen to decrease from 186.3 to 176.2 GPa, withincreasing current density. However, this difference disappears whenthe sample modulus (Esample) is corrected for the anisotropyparameter β using Eq. (1). After this correction the sample modulusvalues were measured to be consistently around 214 GPa, which iscomparable to the modulus value of 219.6 GPa for randomly texturedpolycrystalline Ni [33]. For calculating the values of Esample usingEq. (1), the following values for β were used [28,29], β=0.95, 0.91and 0.89 for Eeff=186.3, 179.5 and 176.2 GPa, corresponding to thethree current densities of 1.5, 2.5 and 5.0 mA cm−2 respectively.

3.4. Residual stress vs. current density

Table 2 shows the evolution of residual stress, calculated usingEq. (2) as a function of the current density. All residual stress valueswere in the range of 74.5 to 135.7 MPa. The thicker deposits (90 μmthickness) were seen to have lower values of residual stress, while thevalues were the highest at 5 mA cm−2 for the 18 μm thick deposit.Two trends are evident from this table. Firstly, the residual stresses areseen to decrease with increasing film thickness – except for the lowestcurrent density of 1.5 mA cm−2, where all three thicknesses havesimilar stress values. And secondly, an increase in current density isseen to induce a steady increase in internal stress of the deposit.

Note that the modulus values for the deposits (Ef) and substrate(Es) used in the calculation of residual stresses in Eq. (2) wereobtained using the indentation results (Esample) described in Eq. (1). Asdescribed in Section 3.3, these values of Esample have been corrected foranisotropy, and as such are comparable to the modulus for randomlytextured polycrystalline Ni.

A more accurate calculation of the residual stresses is also possibleif one uses the effective biaxial modulus (Mf) of the deposit for a giventexture. For example, for the (100) texture (such as the one shown inFig. 2 for the 5.0 mA cm−2 current density),Mf=209.6 GPa [39] usingliterature values for elastic constants of Ni [33], while for the (110)surface the effective biaxial modulus is very close to the randompolycrystalline average [28,29]. However since the texture measure-ments (which require significant experimental resources and ma-chine time) were carried out for only the largest nominal depositthicknesses (90 μm), it was not possible to calculate the effectivebiaxial modulus for all of the textures developed in this work. Insteadthe Esample values obtained from indentation experiments were usedin Eq. (2) for residual stress calculations. This approximation is notexpected to change the main findings of this paper.

4. Discussion

4.1. Crystallographic texture and morphology of electrodeposited Ni

Although deposited at much lower current densities, the mor-phology and texture of the electrodeposited Ni films shown in thiswork share similar features to the Ni films described in literature

Table 2Residual stresses, measured using Eq. (2), and the curvature measurements(±standard deviations) at three different current densities for the three Ni depositthicknesses.

Deposit thickness(μm)

Current density(mA cm−2)

Residual stress(MPa)

Curvature(m)

181.5 76.5 0.058±0.0142.5 131.1 0.192±0.045 135.7 0.187±0.024

691.5 89.7 0.103±0.0052.5 90.3 0.192±0.015 129.6 0.199±0.037

901.5 74.5 0.114±0.0182.5 83.2 0.263±0.0325 94.1 0.221±0.044

produced using higher current densities. The change in the surfacemorphologies of Ni deposits, from a ‘cauliflower’ (or ‘colony’) to a‘truncated pyramidal’ structure, with increase in current density iswell documented in literature [13,14]. Foreign ions, incorporatedduring the deposition process, is thought to block the emerging screwdislocations (which are associated with pyramidal growth), resultingin the truncated structures on the surface. Thus, a higher incorpora-tion of foreign ions into the deposit leads to a decrease in size andeventual disappearance of the pyramids at lower current densities[13]. It is interesting to note here that the transition from the‘cauliflower’ to a ‘truncated pyramidal’ structure is happening at amuch lower current density of 5 mA cm−2 for the 69 μm thick depositsshown in Fig. 1, as compared to 10 mA cm−2 [13] and 30 mA cm−2 [14]for ~100 μm thick deposits reported in literature – a likely consequenceof thedifference in purity of the sulphamate bath solutionsused in thesestudies.

The crystallographic texture of thick Ni deposits, such as the onesstudied in this work, is also known to be related to the grain growth inthese films, which in turn is controlled by the electric field defined bythe plating process [40–42]. Thus at lower current densities theappearance of a fine grained (101) texture is typically associated with‘inhibited’ growth and deposition processes at these conditions, withthe growth mode transitioning to ‘uninhibited’ or ‘free’ growthcharacterized by coarser (100) grains at higher deposition rates(current densities) [3,4,9]. This transition is clearly visible in Fig. 2.Thus, once grains with the favored (100) orientation have occludedmajority of the other grains, columnar grain structures with almostparallel boundaries [40] are produced, such as the ones seen in theOIM cross-section map in Fig. 2 at the highest current density of5 mA cm−2. These results agree well with those of Marquis et al. [9],where a similar transition from (101) to (100) texture, along with amore columnar grain morphology, were observed at 10 mA cm−2

current density for 10 μm thick Ni electrodeposits.

4.2. Microstructure–property relationships

The grainmorphology and texture of the Ni electrodeposits stronglyinfluences their mechanical response during indentation, as shown inFig. 2 and Table 1. Thus an increase in the grain size observed withincreasing current density directly corresponds to a decreasing trend inhardness values, which is a predictable result given the well knowngrain boundary strengthening considerations (such as the Hall-Petcheffect [13,43,44]). Similarly, since the (100) grains have the lowestmodulus values for Ni [28,32], the decrease in the effective modulus(Eeff) with increasing current density is also a direct consequence ofthe change in the grain orientation from (101) to (100). This underlinesthe importance of correcting the sample modulus Esample values for theanisotropyparameterβ, as defined in Eq. (1); otherwise the range of themodulus values can bemisinterpreted as a larger standard deviation orporosity-related changes in the results. As seen from Table 1, once thiscorrection is introduced, themodulus values correspondwell to those ofpolycrystalline Ni.

Interestingly enough, the spread (standard deviation) in the Eeffmodulus measurements for each current density is also around 10% ofthe average (Table 1), which is close to the standard deviationexpected from Vlassak and Nix's model (~10% based on the differencein indentation modulus between the (100) and (101) orientations).This may imply that the Berkovich indenter is probing regions ofvarying texture each time, such that spread is close to the spreadexpected if the indents were evenly distributed between the (100)and (101) orientations. However one needs to carefully correlate eachindent to its underlyingmicrostructure before arriving at any concreteconclusions regarding these correlations. Here we merely point outthe similarity between our experiment and the model developed byVlassak and Nix – a more careful study is currently underway to better

3656 S. Pathak et al. / Surface & Coatings Technology 205 (2011) 3651–3657

understand the influence of the nanoindentation modulus and β tothe exact grain orientation under the indenter.

The indentation modulus values shown in Table 1 compare wellto the values reported in literature using other test methods onNi deposits. For example, a macroscopic sample modulus of 169–179 GPa using axial loading for films deposited at 26.8 mA cm−2 hasbeen reported in [45]. As discussed earlier, at these high currentdensities the Ni film is expected to have a predominantly (100)orientation, and as such these values correspond well to the effectivemodulus value of Eeff=176.2±14 GPa (Table 1) for the (100)oriented 90 μm thick Ni film deposited at 5 mA cm−2. On the otherhand, the hardness (~3.46 GPa) of the 90 μmNi films deposited at thelowest current density of 1.5 mA cm−2, is consistently higher than thevalues reported in literature for Ni electrodeposits prepared in similarsulfamate electrolytes but at higher current densities (for example, aVickers hardness, HV, of 2.3 GPa for 70 μm thick Ni deposits producedat 20 mA cm−2 [46], HV=1.8–2.7 GPa for 30–35 μm thick depositsproduced at 18 mA cm−2 [47], HV~2 GPa for 100 μm thick depositsproduced at 10–200 mA cm−2 [14], HV=2.25 GPa for N100 μm thickdeposits produced at 20 mA cm−2 [3]). This is thought to be a directconsequence of the low grain sizes (~0.8 μm) occurring in the Ni filmsproduced at the unusually low current densities used in this work (ascompared to larger grain sizes of ~1.8 μm reported in [46] for Nideposits produced at 20 mA cm−2), and the resultant grain boundarystrengthening considerations. As such, our results agree well with theHall-Petch type relationship shown by Banovic et al. [13], where asimilar high value (~3 GPa) of Knoop hardness was noted for the finegrain size (b1 μm) Ni deposits plated at low current densities of5–10 mA cm−2.

The low current densities used in this work also result in lowerresidual stress values in the deposited Ni films. As shown in Table 2, theNi films show amonotonic increase in stress with increasing depositionrates (current density) – a trend considered typical for nickel sulfamateplating solutions [1–5,9,10,48,49]. This is generally explained by achange in grain morphology from initially equiaxed small grains tolarger columnar grains, as shown in Fig. 2. Electrodeposition generallyproduces finer grain sizes due to significant impurity incorporation[40]. However the deposition rates of the bath impurities, beingdiffusion limited, do not change with increasing overpotential. As aresult, with increasing overpotential (current density) the impurityconcentration decreases while the Ni deposition rate (which iskinetically limited) increases, leading to grain growth in the deposit[11].

Although columnar grain structures have been noted to evolve athigher current densities leading to a higher residual stress in the Nideposits, the reason behind this stress generation is not clear. Variouspossible mechanisms have been suggested such as a decrease inimpurity concentrations with increasing Ni deposition rate whichcauses the impurity-related compressive stresses to decrease, adecrease in the compressive capillarity stress, as well as adatomtrapping at grain boundaries leading to low adatom mobility [11,49].Of these, the change in adatom concentration has been postulated asthe dominant factor for Ni deposited from sulfamate baths. Indeedduring the deposition, the electrode potential deviates from theequilibrium potential, as determined by Nernst equation, which inturn induces an overpotential η. Since the adatom concentration isdirectly related to η, this results in an increase in the adatomconcentration at higher negative overpotentials (i.e. higher currentdensity) [50]. Small changes in adatom concentration is also known toinduce significant changes in instantaneous stress [51], which wouldexplain the large residual stress changes observed with small currentdensity variations (from 1.5 to 5 mA cm−2) in our deposits. Howevermore work is probably required to verify the above hypothesis. Athicker film is also able to relax more stress through substratebending, as seen in the decreasing trend of residual stress values withincreasing film thicknesses from Table 2.

Thus, the following main trends can be summarized from Table 2 –

for a constantfilm thickness lower residual stress values are obtained fordeposits produced at the lower current density (1.5 mA cm−2), while ata particular current density the thickest film (90 μm) has the loweststress.

4.3. Use of low current densities during electrodeposition of Ni

For the electroplating industry, change in current density is generallyregarded as the most relevant control parameter for electrodeposition,since it determines the process time and the consumed energy. Since theinternal stresses in thick (~microns) deposits are known todecreasewithdecreasing current densities [11], use of very low current densities can beof potential advantage for producing negligible stress values. However,the current work demonstrates that the range in the values of residualstresses (74.5 to 135.7 MPa, Table 2) produced here using unusually lowcurrent densities (1.5–5 mA cm−2) are well within the range of reportedliterature values of internal stresses for Ni electrodeposits produced fromsimilar sulfamate bath compositions (20–90 MPa at 5–25 mA cm−2 for1 μmthickNi deposits [48], ~100 MPa at 50 mA cm−2 for 63–76 μmthickNi deposits [5], 200–250MPa at 20 mA cm−2 for 3 μm thick Ni deposits[52], 150 to 200 MPa at 180 mA cm−2 for 20 μm thick Ni deposits[53,54]). As such, the use of very low current densities for electrodepo-sitionof thickNifilms is not expected to beof any significant advantage inproducing low stress deposits – in fact theymay prove detrimental to theproduction/output speed for the industry.

5. Conclusions

In summary, this work demonstrates the feasibility of producingthick Ni electrodeposits from Ni sulfamate baths using a lower rangeof current densities than typically used in literature. Even at theunusually low current densities used in this work the mechanicalproperties of the electrodeposits are seen to correlate strongly withthe grain morphology and texture of the Ni films. In particular,accounting for the film texture allowed a more accurate calculation ofthe Young's modulus and the residual stresses developed in the films.The higher hardness of the Ni films plated at the lower currentdensities, along with low residual stress values for the larger filmthicknesses makes them an attractive option for the electroforming ofnickel–diamond composites used in high-precision grinding orcutting tools where thicker deposits are desired to keep the diamondsin the nickel matrix. However the residual stress values obtainedusing such low current densities were found to be comparable toliterature values of deposits plated at higher current densities;thus effectively ruling out the attractiveness of employing very lowcurrent densities for producing thick stress-free precise parts usingelectrodeposition.

Acknowledgements

The authors acknowledge helpful discussions with ChristophRudolf, Thomas Sigrist and Marin Pavlovic and financial supportfrom the Swiss Commission for Technology and Innovation (contractnumber 10610.1).

References

[1] N.V. Mandich, D.W. Baudrand, Plat. Surf. Finish. 89 (2002) 68.[2] W.H. Safranek, The Properties of ElectrodepositedMetals and Alloys: A Handbook,

Elsevier Science Ltd, 1974.[3] T. Buchheit, J. Michael, T. Christenson, D. LaVan, S. Leith, Metall. Mater. Trans. A 33

(2002) 539.[4] G.A. Malone, Plat. Surf. Finish. 74 (1987) 50.[5] B.B. Knapp, Planting 58 (1971) 1187.[6] E.C. Lee, J.W. Choi, Surf. Coat. Technol. 148 (2001) 234.[7] W.H. Lee, S.C. Tang, K.C. Chung, Surf. Coat. Technol. 120–121 (1999) 607.[8] Y. Li, H. Jiang, L. Pang, B.A. Wang, X. Liang, Surf. Coat. Technol. 201 (2007) 5925.

3657S. Pathak et al. / Surface & Coatings Technology 205 (2011) 3651–3657

[9] E.A. Marquis, A.A. Talin, J.J. Kelly, S.H. Goods, J.R. Michael, J. Appl. Electrochem. 36(2006) 669.

[10] R.J. Kendrick, in, 1964, pp. 235.[11] S.J. Hearne, J.A. Floro, J. Appl. Phys. 97 (2005)8 014901-014901-014901-014906.[12] J.W. Dini, Electrodeposition: The Materials Science of Coatings and Substrates,

First ed.William Andrew, January 14, 1994.[13] S.W. Banovic, K. Barmak, A.R. Marder, J. Mater. Sci. 33 (1998) 639.[14] A.A. Rasmussen, P. Mller, M.A.J. Somers, Surf. Coat. Technol. 200 (2006) 6037.[15] R. Ghisleni, K. Rzepiejewska-Malyska, L. Philippe, P. Schwaller, J. Michler, Microsc.

Res. Tech. 72 (2009) 242.[16] L. Philippe, P. Schwaller, G. Burki, J. Michler, J. Mater. Res. 23 (2008) 1383.[17] P. Schwaller, L. Philippe, G. Bürki, J. Michler, Proceedings of Viennano'07, March

2007, p. 249.[18] B.L. Adams, Ultramicroscopy 67 (1997) 11.[19] B.L. Adams, S.I. Wright, K. Kunze, Metall. Trans. A 24A (1993) 819.[20] S. Guan, B.J. Nelson, J. Magn. Magn. Mater. 292 (2005) 49.[21] K.P. Mingard, B. Roebuck, E.G. Bennett, M.G. Gee, H. Nordenstrom, G. Sweetman, P.

Chan, Int. J. Refract. Met. Hard Mater. 27 (2009) 213.[22] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564.[23] H. Hertz, Miscellaneous Papers, MacMillan and Co. Ltd., New York, 1896.[24] K.L. Johnson, Contact Mechanics, Cambridge University Press, Cambridge, 1987.[25] S. Pathak, S.R. Kalidindi, C. Klemenz, N. Orlovskaya, J. Eur. Ceram. Soc. 28 (2008)

2213.[26] S. Pathak, D. Stojakovic, S.R. Kalidindi, Acta Mater. 57 (2009) 3020.[27] J.J. Vlassak, M. Ciavarella, J.R. Barber, X. Wang, J. Mech. Phys. Solids 51 (2003)

1701.[28] J.J. Vlassak, W.D. Nix, Philos. Mag. A Phys. Condensed Matter Defects Mech.

Properties 67 (1993) 1045.[29] J.J. Vlassak, W.D. Nix, J. Mech. Phys. Solids 42 (1994) 1223.[30] J.R. Willis, J. Mech. Phys. Solids 14 (1966) 163.[31] G. Simmons, H. Wang, Single Crystal Elastic Constants and Calculated Aggregate

Properties. A Handbook, 2nd edition ed.The MIT Press, 1971.

[32] W.F. Hosford, The Mechanics of Crystals and Textured Polycrystals, OxfordUniversity Press, USA, 1993.

[33] Z. Hashin, S. Shtrikman, J. Mech. Phys. Solids 10 (1962) 343.[34] M. Ohring, Materials Science of Thin Films, Second edition, Academic Press, San

Diego, 2002, p. 711.[35] G.G. Stoney, Proc. R. Soc. London Ser. A 82 (1909) 172.[36] T.W. Clyne, Key Eng. Mater. 116–117 (1996) 307.[37] T.W. Clyne, S.C. Gill, J. Therm. Spray Technol. 5 (1996) 401.[38] H.E. Boyer, T.L. Gall, Metals Handbook, American Society for Metals, Materials

Park, OH, 1985.[39] L.B. Freund, S. Suresh, Thin Film Materials: Stress, Defect Formation and Surface

Evolution, Cambridge University Press, 2004.[40] C.V. Thompson, Annu. Rev. Mater. Sci. 30 (2000) 159.[41] C.V. Thompson, R. Carel, J. Mech. Phys. Solids 44 (1996) 657.[42] C.V. Thompson, R. Carel, in, Trans Tech Publications, Switzerland, 1996, pp. 83.[43] E.O. Hall, Proc. Phys. Soc. Sect. B 64 (1951) 747.[44] N.J. Petch, Iron Steel Inst. J. 174 (1953) 25.[45] J.W. Dini, H.R. Johnson, Surf. Technol. 4 (1976) 217.[46] C.S. Lin, P.C. Hsu, L. Chang, C.H. Chen, J. Appl. Electrochem. 31 (2001) 925.[47] F. Ebrahimi, G.R. Bourne, M.S. Kelly, T.E. Matthews, Nanostruct. Mater. 11 (1999)

343.[48] J.J. Kelly, S.H. Goods, A.A. Talin, J.T. Hachman, J. Electrochem. Soc. 153 (2006) C318.[49] A. Bhandari, S.J. Hearne, B.W. Sheldon, S.K. Soni, J. Electrochem. Soc. 156 (2009)

D279.[50] P.D.E. Budevski, P.D.G. Staikov, P.D.W.J. Lorenz, Electrochemical Phase Formation

and Growth, 2007, p. 9.[51] C. Friesen, C.V. Thompson, Phys. Rev. Lett. 89 (2002) 1261031.[52] Y.C. Zhou, Y.P. Jiang, Y. Pan, Adv. Mater. Res. 9 (2005) 21.[53] Y. Tsuru, M. Nomura, F.R. Foulkes, J. Appl. Electrochem. 32 (2002) 629.[54] Y. Tsuru, M. Nomura, F.R. Foulkes, J. Appl. Electrochem. 30 (2000) 231.