Cavitation erosion of electroplated nickel composite coatings
Thomas Lampke, Dagmar Dietrich ⁎, Anette Leopold, Gert Alisch, Bernhard Wielage
Institute of Composite Materials and Surface Technology, Chemnitz University of Technology, Germany
Received 15 November 2007; accepted in revised form 10 February 2008Available online 16 February 2008
Abstract
The cavitational wear resistance of electroplated nickel composite layers was tested following ASTM G32. Particles of different hardness(titania and silicon carbide) and different sizes from micro-scale to nano-scale were incorporated up to 30 vol.% into a nickel matrix. Martenshardness is improved by grain refinement via particle incorporation. Compared to pure electroplated nickel films the composite layersstrengthened by submicro-scale silicon carbide particles exhibit a decreased mass loss of one order of magnitude after 8 h testing time.Remarkably, layers with nano-scaled titania particles show a similar performance.
Nickel and nickel dispersion coatings are used in a multitudeof applications where corrosion and wear resistance isrequested. Dispersion films combine the ductility of the metalmatrix with the hardness of incorporated non-metallic, mostlyceramic particles like alumina, titania or silicon carbide. Thetribological properties of dispersion films are mainly deter-mined by the content, size and properties of the particlesand their dispersion. For automotive engines, nickel compositelayers with silicon carbide respectively lubricant particles onaluminium are state-of-the-art [1] and, for example, the long-lifeperformance of Formula 1 race engines are enabled by suchcomposite layers [2].
The application-oriented development of such layers neces-sitates complex research about optimized deposition conditionsyielding microstructural properties of the composites which
⁎ Corresponding author. Chemnitz University of Technology, Institute ofComposite Materials and Surface Technology, D-09107 Chemnitz, Germany.Tel.: +49 371 531 35392; fax: +49 371 531 23819.
correlate with mechanical materials characteristics like hard-ness, adherence, ductility, strength and wear resistance [3–5].Different indentation tests, e.g. Vickers, Martens, and thescratch test, are well established techniques for the evaluation oflayers and composites. Nevertheless, these techniques exertonly static stress on the materials. Normally, structural com-ponents are subjected to complex dynamical loads wherebyfatigue cracks, fretting damage and corrosion may occur. Apartfrom failures resulting from vibration stress, cavitation occurs tovarious degrees in fluid-handling equipment. Propellers, turbinepumps and pipelines are important parts of manufacturingindustries, where nickel composites are interesting corrosion-and erosion-resistant layers. Nevertheless, little informationappears pertaining to the effectiveness to resist cavitationalstress. Despite of theoretical and experimental efforts, no fullysatisfactory solution has been found so far setting up a rela-tionship between materials properties an their cavitation resis-tance. Solely one paper could be found regarding to dispersioncoatings (electroless Ni–P–SiC) and their cavitation exposure[6].
The cavitation test according to ASTM G32 [8] is one ofthe established tests that have been developed for evaluatingcavitation resistance. Under the influence of ultrasound, the
Fig. 1. Well-dispersed particles under ultrasound conditions (lower part fromsubstrate up to the markers Nb) and nano-particle agglomeration under silentconditions (upper part).
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repetitive formation and collapse of cavities in a liquid gen-erates shock waves at a regular frequency and micro-jets. Thekinetic energy is transformed into heat and mechanical energy.“Hot spots” with temperatures above 5000 K and shock wavescausing loads in the range of GPa can be generated [9].Materials which are subjected to a combination of impact andfatigue stresses undergo elastic and/or plastic deformation at thesurface. This leads to fatigue, fracture and material loss.Depending on the mechanical properties a network of crackscan develop. When the cracks propagate and join, small par-ticles break loose, leaving behind a pitted surface [10].
This paper deals with the erosion resistance of nickel dis-persion layers under cavitational load. The wear resistance ofcomposites can be enhanced by strain hardening via formation offine grains and incorporation of dispersed particles. Thereforethe examination of nickel composite layers was focused on thecorrelation between the microstructure and the materials be-haviour under dynamical load. Inorganic particles from micro-to nano-scale were incorporated to study particle size effects. Inaddition to SiC particles TiO2 was used for differentiation ofparticle features like hardness or surface properties. TiO2 waschosen by others due to its outstanding potential as photo-catalyzer to be used as functional material [11]. SEM studies onparticle distribution, matrix grain formation and crack propaga-tion were compared to cavitational mass loss and erosion rates ofnickel and nickel composite layers. Correlations to Martenshardness and the ratio of plastic and elastic work acquired fromthe instrumented indentation test were found and discussed.
2. Experimental
2.1. Deposition conditions
Nickel composite layers were electroplated on polished disks(steel 1.0503-C45, 40 mm diameter, 10 mm height) in Wattselectrolyte (250 g/l NiSO4
. 7H2O, 30 g/l NiCl2. 6H2O, 40 g/l
H3BO3, 0.3 g/l NaC12H25SO4). A film thickness of 100 μm wasreached after 4 hours under following deposition conditions:cathode current density 400 A/m2, temperature 50–55 °C, pH-value 3.2–4.2, stirring at 700 rpm. A solids concentration of20 g/l in the electrolyte was used for particle incorporation. Inorder to avoid agglomeration of nano-scaled particles ultra-sound was applied during electrodeposition by means of aheated ultrasonic bath with a frequency of 35 kHz (Sonorexsuper, Bandelin). All coatings were obtained at same conditionsexcept for different materials (titanium oxide and silicon car-bide) and particle sizes.
Table 1Producer specification and phase analysis of the incorporated particles
Primarily, particles for incorporation were selected in view oftheir size effect on particular matrix formation. TiO2 and SiCwere provided according to producer's specifications in micronsize (5–20 μm respectively 5 μm), in submicron size (280 nmrespectively 550 nm), and TiO2 additionally in nano-scale size(21 nm). Table 1 summarizes the producer's size specificationand the results of our X-ray diffraction phase analysis. XRDwas done in a XRD diffractometer D5000 (Siemens). Scanningelectron microscopic images of the powders revealed the di-verse particle shape, which is fractured in the case of the largerparticles and rounded in the case of the submicron- and nano-scaled particles. Aside from chemical composition the hardnessof the particles is another differing feature. The Mohs hardnessof TiO2 between 5.5 for rutile and 6.5 for anatase is comparableto nickel (about 5); the Mohs hardness of SiC is remarkablyhigher (9.6).
2.3. Microstructure investigations and wear tests
Samples for scanning electron microscopy (SEM LEO1455VP) were pressure-embedded, mechanically grinded andoxide-polished. SEM studies were combined with energydispersive X-ray spectroscopy (EDX detector, Kevex) for es-timating the particle content in the composite films. Back-scattered electron beam diffraction (EBSD) data were collectedin a Nova nano SEM 2000 (FEI) equipped with a Nordlys IIelectron backscattering diffraction detector (Oxford Instru-ments) for studying textures and measuring grain sizes. Aninstrumented indentation tester (FISCHERSCOPE HM100)was employed for measuring hardness and materials parametersaccording to ISO 14577 [7]. Martens hardness (HM 0.5) and theplastic part of the indentation work were determined for allcomposite layers.
The cavitation erosion resistance was evaluated according toASTM G32-92 (vibratory frequency 20 kHz, peak-to-peakamplitude 65 μm, temperature 25 °C, normal pressure) using an
Table 2Mean grain intercepts normal to and in film growth direction, aspect ratios ofnickel grains derived from intercepts
Sample Intercept normal togrowth direction
Intercept in growthdirection
Aspectratio
[μm] [μm]
Standard Watts Ni 0.49±0.47 0.68±0.73 1.4Ultrasonic impact 0.38±0.25 0.56±0.60 1.5Ni+21 nm TiO2 0.27±0.23 0.28±0.20 1.0Ni+280 nm TiO2 0.36±0.30 0.57±0.87 1.6Ni+5–20 μm TiO2 0.26±0.19 0.33±0.35 1.3Ni+550 nm SiC 0.28±0.20 0.34±0.38 1.2Ni+5 μm SiC 0.51±0.26 0.66±0.92 1.3
ultrasonic device (UIP 250, Dr. Hielscher Ldt.). Apart from theincreased amplitude a commonly used modification was intro-duced by using an indirect test arrangement [8,12]. The testspecimen was situated at a distance of 0.5 mm to the oppositevibrating sonotrode end face. The test was performed in distilledwater on layer surfaces defined by fine grinding (grid 600).The total exposure time was 8 h; the incremental mass loss wasmeasured hourly.
3. Results
3.1. Microstructure
The electroplated composite layers show homogeneous par-ticle dispersion under constant deposition conditions. This wasreached for submicron and nano particles by application ofultrasound to the electrolytic bath. Fig. 1 reveals that the nano-particles are homogenously incorporated as long as the ultra-sound acts on the electrolyte and they immediately agglomeratewhen the exposure was stopped. The particle concentrationincreases with increasing concentration of the solids content inthe electrolyte and increasing particle size. Under similardeposition conditions a particle concentration around twice asmuch as for titania can be reached for silicon carbide.
The film growth of the nickel matrix starts with an initiallayer of about 1 μm thickness with small unoriented crystallitesfollowed by a subsequent columnar growth with fibre texture.The crystal size and texture is influenced by ultrasound ex-posure as well as the nature and size of the incorporated par-ticles. As mean grain size parameters the mean linear interceptsin film growth direction as well as normal to film growthdirection were derived from EBSD quality maps. Values rangebetween 200 and 900 nm. The similar and fairly high standarddeviations are due to a broad grain size distribution. Because ofthe mainly columnar nickel growth the aspect ratio was cal-culated as the quotient of the mean intercepts. Table 2 resumesthese grain size parameters of the nickel and nickel composites.
Fig. 2. Different grain growth of nickel films, orientation maps in randomizedgrey scale, high angle boundaries in black, low angle boundaries (2–10°) ingrey, representative pole figure inserted. a) Nickel deposited under standardconditions. b) Nickel deposited under ultrasound conditions. c) Nickel with TiO2
nanoparticles.
Fig. 5. The influence of particle content on plastic deformation of the composite.
Fig. 4. The influence of particle content on Martens hardness of the composite.
Fig. 3. Different grain growth around micron particles, high angle boundaries inblack, low angle boundaries (2–10°) in grey, orientation maps in randomizedgrey scale, representative pole figure inserted. a) Nickel with TiO2 microparticles. b) Nickel with SiC micro particles.
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In conjunction with the EBSD derived orientation maps(Fig. 2 a–c) they show that ultrasound provokes a higher aspectratio of the columnar grains and particle incorporation causes atransition from columnar to more granular crystal growth.
Pure electroplated nickel frequently shows b211N fibre tex-ture in growth direction. The co-deposition of particles as wellas the application of ultrasound to the electrolyte during theelectroplating process influences the formation of textures.They change to a b110N fibre texture for deposition underultrasonic exposure and mostly to a b100N texture in the case ofparticle incorporation. Remarkably, micro-particles of titaniaare surrounded by individual layers of nickel crystals startingsimilar to the initial layer on the cathode surface (Fig. 3a);consequently no texture formation can be observed in suchcomposite films. In contrast, large columnar nickel crystalsgrow undisturbed between silicon carbide particles of micronsize and a fibre texture develops. We attribute the differentgrowth forms to the isolating character of titania and the semi-conducting character of silicon carbide. Nickel crystals grow asfar as they are terminated by a SiC particle and only on the topof these particles the electroplating starts with a new fine crys-talline initial layer (Fig. 3b).
The particle concentration in the deposits is outlined in termsof percent by volume. It is influenced by the particle size, the
particle composition and the particle load in the electrolytic bathwhich was always 20 g/l. We measured the particle content ofall samples before indentation and wear investigations byEDXS, since it can not be derived reliably from depositionconditions. The data are calculated from the Ti/Ni ratio and Si/Ni ratio, respectively, assuming the stoichiometric composition,since light elements, especially carbon and oxygen hardly canbe analysed quantitatively by EDXS. The microstructure resultssummarized in this section have been reported in detail else-where [13].
3.2. Hardness
Compared to composite layers, the pure nickel layers gen-erally show the lowest values of the Martens hardness whetherstandard (larger crystals) or ultrasonically aided (smaller crys-tals) electroplating. Strikingly, the Martens hardness increaseswith increasing particle content in the composite and decreasingparticle size (Fig. 4). In this regard the behaviour of nano-scaledtitania-strengthened composites is notable. The diagram showsthis general influence, although a direct correlation was notfound. Nevertheless a tendency can be drawn taking into ac-count the pointed lines. Similarly the relationship between par-ticle content, particle size and the plastic ratio of indentationwork (Fig. 5) was illustrated.
The instrumented indentation test provides the advantage ofmeasuring supplementary mechanical properties of the materialapart from Martens hardness, e.g. the plastic fraction of
Fig. 6. Relationship between mass loss and exposure time during cavitationalwear testing.
indentation work. This is the quantity of formability of amaterial and allows an estimation of the probability of crackformation. Likewise the Martens hardness the fraction of plasticdeformation mostly decreases with increasing particle concen-tration and decreasing particle size and both are associated withdecreasing nickel grain size.
3.3. Cavitational wear
Fig. 6 summarizes the mass loss by cavitational wear duringa test period of eight ours. The mass loss of pure nickel notice-ably decreases by incorporation of particles. The lowest massloss is reached by incorporation of sub-micro silicon carbideparticles (550 nm) whereas the effect of sub-micro titaniaparticles (280 nm) is marginal. Nickel with incorporated micro-scaled titania particles increases compared to nickel withoutparticles.
A direct comparison of cavitational resistance between dif-ferent materials permits the mean depth of erosion and theerosion rate. The mean depth of erosion is determined as theaverage thickness of material removed from a specified surfacearea and was calculated by dividing the measured mass loss bythe density of bulk nickel and the damaged surface area. Basedon the mean depth of erosion the corresponding erosion rate was
Fig. 7. Mean erosion rate as a function of exposure time.
calculated (Fig. 7). However, the mean erosion rate is based onthe assumption of one-dimensional material removal, whilecavitational wear is a three-dimensional process leavinga sponge like structure or a so-called ‘honeycomb’ appearance.For that reason mass loss and erosion rate results will be dis-cussed in terms of actual damage of the composite layers whichwere observed by scanning electron imaging on treated surfacesand cross sections.
Typically three stages can be distinguished in the incrementalincrease of the mean erosion rate during the first hours. Duringthe initial or incubation stage, no material loss is observed. Theincubation period, i.e. the intercept on the time axis of anextension of the maximum slope of the erosion-time curve,ranges from a few minutes for pure nickel and nickel withincorporated micron-scaled titania to 3 h for silicon carbide-strengthened layers. In the second phase, called accelerationstage, mass loss begins expressed by a rapid rise of the erosionrate. Nickel with incorporated micron-scaled titania shows thehighest increase of the erosion rate. The third phase, the maxi-mum rate stage, is characterised by constant erosion expressedin terms of a constant erosion rate. The maximum rate of erosionproposed by ASTM as a representative number for comparisonof cavitation tests under the same conditions will be discussed inrelation to the Martens hardness of the tested coatings (Fig. 13).
4. Discussion
Contrary to other wear types, cavitational wear cannot berealized at the beginning. Invisible changes like lattice defectsand distortions occur during the so-called incubation period.The variety of cyclic events accumulates stress and initiatesdislocations, causes work hardening, fatigue and cavitation pits.The surface region will be deformed, ductile fracture failure orbrittle fracture failure occurs depending on materials behaviour.A successful prediction of cavitation erosion still remains oneof the goals in this field of research since the process of pitformation is very complex. In [14] the pit formation is sug-gested in the following way: the collapse of the cavitation cloudcauses a shock wave that spreads in the fluid. The magnitude of
Fig. 8. Eroded surface of a nickel layer after 8 h testing time.
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the shock wave is damped by the fluid. As the shock wavereaches the solid surface, single bubbles on it begin to oscillateand a micro-jet phenomenon can occur. The damage by singlepits is caused by the high velocity liquid jet impact to the solidsurface. Under cyclic loading of the imploding cavitationalbubbles the propagation of initiated cracks forms a network.
Fig. 10. Typical pitting of nano-scaled titania-strengthened nickel after 8 hours.
Fig. 9. Cross sections of eroded nickel titania composite layers after 8 h testingtime. a) Pure nickel. b) Ni composite with 5 μm TiO2. c) Ni composite with280 nm TiO2.
When the cracks propagate they provoke discrete regions nearthe surface to be ejected as debris. The eroded surface shows a‘honeycomb’ appearance; for example the pure nickel layer
Fig. 11. Typical surface erosion of micro-scaled titania nickel composite after8 h. a) Periphery of erosion. b) Centre of erosion.
Fig. 13. The correlation of Martens hardness with maximum erosion rate.
Fig. 12. Nickel silicon carbide dispersion layers after 8 h testing time. a) Typicalsurface erosion of a Ni composite with 5 μm SiC. b) Cross section of a Nicomposite with 5 μm SiC.
(Fig. 8) shows a severely eroded surface mainly by ductilefracture failures after eight hours of cavitation testing.
Cracks develop involving progressive loss of material at thesurface. The crack propagation seems to follow the boundariesof columnar nickel grains or the interfaces between incorpo-rated particles and surrounding nickel grains. Severe erosion iscaused by the coherent failure between large columnar grainswith high aspect ratio in the case of pure nickel to be shown bythe cross section (Fig. 9a). The nickel composite with mico-scaled titania particles fails due to the low adherence of theparticles and subsequent erosion is accelerated due to largepitting (Fig. 9b). Crack propagation can be stopped by finergrains, lower aspect ratios of the crystals and by well adherentparticles. For example incorporated nano-scaled titania particleswith good adherence retard the erosion by grain refinement.Smaller fragments of material are removed which reveal thecross sections (Fig. 9c) as well as the appearance of the sur-face (Fig. 10). This results in a halved mass loss after 8 h testingtime. The crack propagation in the metallic matrix is hinderedby incorporated particles since a portion of the crack propa-gation energy is consumed to overcome the obstacles by energydissipation. This gives rise to a rather homogenous material
removal. Not only erosion of such composites is decelerated butalso the incubation time is improved and gives rise to a betterwear resistance.
Scanning electron microscopic imaging of the periphery of theerosion zone of nickel with micro-scaled titania shows the lowadherence of these particles (Fig. 11a). They can be ejected easilyby increasing cavitational stress which occurred in the centre ofthe affected region (Fig. 11b). Erosion then proceeds in the holesof some microns and quickly degrades the surrounding nickelmatrix. The collapsing cavitational bubbles act on a steadilyincreasing surface and the erosion rate increases dramatically.
Silicon carbide nickel composite layers show a remarkablyimproved wear resistance even in the case of micro-scaledparticles. The erosion rate after eight hours testing time isdecreased by one order of magnitude compared to the purenickel matrix. Only small pitting and shallow cracks can beobserved after eight hours testing time for composites withmicro-particles of SiC (Fig. 12a and b).
The better wear resistance of SiC strengthened nickel is morepronounced for submicro-scaled particles which is similar torecently published results for electroless nickel composites. Theincorporation of SiC particles greatly reduced the occurrence ofpitting in electroless Ni–P coatings. Nano-particles provided thegreatest cavitation erosion resistance, and appeared to inhibitthe onset of erosion damage [4]. Moreover, this is in accordancewith our results, especially the remarkable performance of thenano titania composites.
The correlation between Martens hardness of the compositesand the maximum rate of erosion after 8 h testing time isdemonstrated in Fig. 13. An exception is the dramaticallyincreased maximum rate of erosion of the composite with mico-scaled titania particles compared to pure nickel despite a com-parable Martens hardness. As discussed above the low cavi-tational resistance is primary due to the poor adherence of thetitania particles which left holes of some microns acting asdegradation centres. Sufficient particle adherence is a require-ment for dispersion hardening effects of composite materialsachieved by particle incorporation. Hereby cavitational wearresistance remarkably improves. This is in agreement withestablished models which predict that the energy absorbed bythe plastic deformation of the material depends on the hardnessof the material [15]. Best hardening and wear resistance was
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reached in this study by incorporating submicron silicon car-bide. The composite combines good particle adherence, appro-priate particle spacing and matrix grain refinement, whichresults in the highest composite strengthening. Consequentlythe good performance of nickel composites with submicron ornano-scaled titania particles is not unexpected. This effect ismore striking for nano-scaled titania-strengthened nickel wherethe erosion damage occurs evenly on the surface and has nopreferential locations.
5. Summary
From the investigations on electroplated nickel and nickelcomposite layers the following aspects can be stated. In purenickel films fatigue cracks frequently develop and propagatealong boundaries of columnar nickel grains during cavitationalstress. Discrete regions near the surface are ejected mainly byductile fracture failures. By incorporation of particles like titaniaor silicon carbide fatigue crack propagation can be hindered andthe erosion damage occurs evenly on the surface. The cavi-tational wear resistance of nickel composites can be improvedif particle-matrix bonding is sufficient.
If particle-matrix bonding is insufficient the particles willbe pulled out of the matrix by persistent cavitational stress.Adhesive failures give rise to subsequent accelerated erosion ofthe matrix due to large pitting. A sponge-like structure developswith partially subsurface erosions. Incorporation of submicro-scaled particles causes granular grains and grain refinement.Nano-scaled particles cause not only the smallest mean grainsize but also a narrow grain size distribution with the con-sequence of improved hardness. If particle-matrix bonding isadequate the maximum erosion rate due to cavitational wear ismainly correlated with the Martens hardness of the composite.The results show that properly selection of particles (size, type)and identification of plating parameters leading to good in-terface bonding conditions is the key to improve compositematerials resistance against cavitation attack.
Acknowledgements
We greatly acknowledge the cooperation with the Dr.Hielscher Ltd., Teltow, providing the sonotrode used forcavitational investigations. Additionally, we like to thank Dr.H. Podlesak and G. Fritsche, Institute of Composite Materials
and Surface Technology, for TEM and XRD investigationsand helpful discussions. Prof. M. Hietschold, Solid SurfacesAnalysis, Institute of Physics and the technical assistance of hisstaff are gratefully acknowledged.
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Dr. Thomas Lampke was born in 1968 and studied mechanical engineering inBremen and Materials Science in Chemnitz and is the chief engineer of theinstitute.
Dr. Dagmar Dietrich was born in 1953 and studied physics in Chemnitz and isresponsible for the structural characterisation of materials.
Dipl.-Ing. Anette Leopold was born in 1972 and studied materials science inChemnitz. She is production engineer at AMD Dresden.
Dr. Gert Alisch was born in 1948 and studied Materials Engineering inChemnitz and is responsible for anodizing light metals and mechanicalcharacterization of materials.
Prof. Dr. Habil. BernhardWielage was born in 1946 and studied engineering andmaterials science in Hannover. He is dean of the Faculty of MechanicalEngineering and head of the Institute of Composite Materials at ChemnitzUniversity of Technology.
