Current Opinion in Solid State and Materials Science 9 (2005) 211–218

Role of metal interlayers in the infiltration of metal–ceramic composites

Carlos A. Leon-Patino a, Robin A.L. Drew b,*

a Instituto de Investigaciones Metalurgicas, Universidad Michoacana de San Nicolas de Hidalgo, Apdo. Postal 888, c.p. 58000 Morelia, Mich., Mexicob Metals and Materials Engineering, McGill University, 3610 University Street, Montreal, Canada H3A 2B2

Received 20 March 2006; accepted 4 April 2006

Abstract

Interlayers using electroless plating technologies have been demonstrated for coating ceramics in the production of activated powders.Of particular note is the ability to enhance metal infiltration during the production of metal–matrix composites. However, recent break-throughs have highlighted the diversity of the technique using different chemistries that are adaptable to producing interlayers onhydrides and other powders.� 2006 Elsevier Ltd. All rights reserved.

Keywords: Interlayers; Ni coating; Electroless plating; Infiltration; Metal–ceramic composite

1. Introduction

Metal–ceramic interlayers play an important role in thefabrication of metal–ceramic and metal–matrix compos-ites. They may be employed as coatings to ceramic particlesto activate their normally inert surface or act as protectionfrom reaction with liquid metal. Moreover, metal coatingscan promote infiltration of liquid metal into ceramicpreforms during the fabrication of MMCs. This paper dis-cusses the use of such interlayers in this important area ofmaterials engineering.

Composite materials are a synthetic combination ofmaterials to achieve a particular function. They haveemerged because of the need for improved properties ofmaterials and metal–matrix composites constitute one ofthe most important groups of composites.

The purpose of the matrix in a composite is to incorpo-rate reinforcing particles and take up and distribute exter-nal loads within the material bulk. In addition, the matricesprotect the reinforcing phase against external effects:mechanical damage, erosion and corrosion caused by the

1359-0286/$ - see front matter � 2006 Elsevier Ltd. All rights reserved.

doi:10.1016/j.cossms.2006.04.006

* Corresponding author. Tel.: +1 514 398 1773; fax: +1 514 398 4492.E-mail address: robin.drew@mcgill.ca (R.A.L. Drew).

surrounding medium or reinforcement–reinforcement con-tact. The most common alloys employed as matrices arealuminum, titanium and magnesium, as well as some inter-metallic compounds [1–3]. Titanium has been extensivelystudied for higher-temperature applications. Nickel andnickel-based superalloys are also possible candidates forhigh-temperature applications when weight limitation isnot a primary objective [4]. However, aluminum is the moststudied metal–matrix, especially for applications at temper-atures below 450 �C, owing to its low cost and greatestpotential in terms of the strength to density ratio. Further-more, aluminum matrix composites have high electricaland thermal conductivity.

The chemical composition of the matrix alloy is veryimportant in solving the problem of matrix-reinforcementcompatibility. In the case of the liquid-metal processingfor composites production, both cast and wrought alloyshave been used. The choice of the alloy can depend heavilyon the final property requirements and the processing routeinvolved; for making net-shape components, cast alloys aremostly used. On the other hand, for making semi finishedwrought products (e.g., billets, rods, etc.) wrought alloyscan be used [5,6].

In most cases industrial ceramic powders are used forcomposite fabrication, i.e., oxides, carbides, borides and

212 C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218

nitrides. The sizes of particles range from one to tens ofmicrometers, depending on the purpose of the compositeand the matrix alloy used. For strengthening aluminumalloys, the most widely used powders are SiC and Al2O3,and for heat-resistant nickel alloys, powders of ThO2,ZrO2 and Y2O3 are normally employed [7,8]. However,the choice of reinforcement is dictated by factors like par-ticle shape, method of composite manufacture, cost andapplication. With these considerations in mind, the tworeinforcements receiving the most attention are SiC andAl2O3.

Liquid-metal infiltration may be used for the fabricationof metal–matrix composites. The method comprises of aprocess by which the gases within a porous body are pro-gressively replaced by a low viscosity liquid melt, thuseliminating the solid/vapor interface in the system. Liquid-metal infiltration methods for fabricating composite mate-rials are favorable for two reasons: firstly, small diameterparticles or whiskers can barely be introduced into a solidmatrix; secondly, liquid infiltration methods are relativelysimple [9,10]. Normally, the process is aided by the applica-tion of an external force, such as pressure or vacuum.

Liquid metal can be driven into a porous preform sim-ply by capillarity. The preform is an agglomerate or a com-pact of the reinforcement with essentially the same shape asthe final MMC component. Pressureless melt infiltrationrefers to liquid metal infiltration without an applied pres-sure. The elimination of the need for application of pres-sure greatly simplifies the process. The attraction of thismethod is that the preform may be conveniently shapedor machined to the desired shape of the composite, andnear-net-shape component after liquid metal infiltration.Furthermore, depending on the composite system, themethod can be rapid compared to powder metallurgymethods [11–14].

Spontaneous infiltration can be achieved by improvingwetting between the reinforcements and the molten alloysthrough certain chemical reactions. To improve wettingof ceramic particles by liquid metals, one approach hasbeen to apply a metal coating to ceramics, which essentiallyincreases the overall surface energy of the solid (cSV),thereby promoting wetting by the liquid metal [15]. Severalinvestigators have attempted to fabricate MMCs withmetal-coated reinforcements by liquid infiltration and pow-der metallurgy techniques. The ceramics employed wereboth fibers and particulates of SiC, Al2O3, TiB2 and graph-ite, while the coatings used are copper, nickel, silver andtitanium [16–18].

Cheng and Chung [19] fabricated, under a nitrogenatmosphere, Al–matrix SiC composites by both pressure-less and pressure infiltration of liquid Al–Mg and Al–Si–Mg alloys at 830–950 �C into preforms of Ni-coated SiCwhiskers. The composite exhibited slightly lower tensilestrength and modulus, but slightly higher coefficient ofthermal expansion than the corresponding composite madeby pressure infiltration. These differences are attributed tothe higher porosity of pressureless infiltrated samples.

Chung and Lin [20,21] measured the strength of Ni-coatedSiCP reinforced aluminum composites fabricated by vac-uum infiltration of Al alloys at 650–800 �C. It was foundthat the bend strength of the composites decreased as thevolume fraction of SiC increased; the strength varied fromabout 270 MPa for the composites containing 5 vol.% SiCto 120 MPa for composites containing 48 vol.% ceramic.Extensive segregation of nickel was observed in the matrix.Furthermore, Al–Ni phases such as NiAl3 were found as aresult of nickel dissolution.

Asthana and Rohatgi [22] attributed a partial but bene-ficial effect of Ni coating on the penetration length of Ni-coated SiC particulate preforms with 2014 Al and Al–Sialloys. They reported undissolved nickel enveloping SiCparticles after infiltration. Murty et al. [23] suggested thatCu and Ni coatings do not significantly influence the infil-tration behavior of SiC with pure aluminum at low preheattemperatures (500 �C). However at higher preheat temper-ature (650 �C), infiltration of SiC was enhanced in the pres-ence of coatings. Furthermore, no major improvement wasobserved in the infiltration kinetics with Ni and Cu coat-ings. A clear trend of increasing hardness from theuncoated, Cu-coated through Ni-coated SiC was observed.On the other hand, results on the infiltration of 2014 Alinto porous compacts of plain and Cu-coated SiC particlesby pressure infiltration indicated a positive effect of coppercoating on the threshold pressure as a result of improve-ments in wetting [24]. Greater infiltration lengths wereachieved in coated samples compared to uncoated onesunder identical processing conditions. Moreover, coarserSiC led to greater total lengths than fine SiC. Bhanuprasadet al. [25] found an improvement in strength from 85 MPato 105 MPa of composites reinforced with Cu-coated SiCparticles compared to uncoated Al/SiC composites,although with a decrease of 8% in elongation.

Hojo et al. [26,27] fabricated multiphase composites byinfiltrating molten copper and nickel into porous sinteredbodies consisting of agglomerates of Ni-coated Al2O3 par-ticles. Molten copper and nickel were infiltrated into per-forms at 1100 �C and 1400 �C, respectively, giving thedense composites of Al2O3–Ni and Al2O3–Ni–Cu. The metal-infiltrated composites were dense regardless of the size ofAl2O3 particles and the relative density reached 95–99%.The easy penetration of molten metals into pores is dueto the metal interlayer on Al2O3 particles. Alternatively,Cu-coated SiC particles are used for processing dispersionstrengthened copper matrix composites by hot pressing,producing materials with low porosity and high particulatevolume fraction. The mechanical properties of those com-posites were enhanced about three times over those of purecopper, yet ductility is partly maintained when the speci-mens are fully dense [28,29].

2. Coating and infiltration of SiC particles

Autocatalytic nickel plating was performed on two sam-ples of a-SiC powders (97% min, Norton Silicon Carbide)

C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218 213

with average particle sizes of 49 and 78 lm. The coatedpowders were infiltrated with commercially pure aluminum(Alcan Aluminum Ltd.) containing minor impurities of Siand Fe; the alloy has a yield strength of 30 MPa and a ten-sile strength of 70 MPa. To ensure adequate coating adhe-sion and initiate deposition on ceramics, every particle wascleaned, microetched, and surface catalyzed prior to plat-ing [30]. The Pd-activation procedure made the particlessuitable for electroless plating due to the presence of chem-ically deposited Pd nuclei that initiate nickel deposition.Autocatalytic nickel deposition was performed usingsodium hypophosphite as the reducing agent, giving elec-troless nickel deposits along with a simultaneous deposi-tion of phosphorus from the electroless bath. Therefore,the electroless nickel deposit was a Ni–P alloy, as con-firmed by EDS.

All the coated-powders were tested and infiltrated withmolten aluminum using a vacuum infiltration technique.The powders were packed into quartz tubes in the as-coated conditions and the matrix loaded into a clay cruci-ble and placed inside an electrical furnace of the infiltrationsetup [31]. The powders were poured into the 0.8 cm inter-nal diameter tube in batches of 0.8–1 g until a compactheight of 6 cm was reached. Between batches, the particu-lates were packed with a tamping device to ensure noimperfections or discontinuities due to powder layeringduring compacting that might influence the infiltration pro-cess. Immediately the prepacked reinforcement wasinserted into the melt, a vacuum pump fitted to the tubewith a rubber seal mechanism was turned on and the infil-tration time was considered to start. After completeinfiltration or a predetermined infiltration time, the as-infil-trated sample was removed from the melt, the pump wasturned off and the tube was vented through a three-wayvalve. Subsequently, the composite was allowed to cooland solidify in ambient air. Sample cross sections were pre-pared for SEM, EDS and XRD analysis. The density,hardness, elastic modulus, flexural and tensile strengthof the composites were evaluated through standardprocedures.

Previous work performed on the effect of electrolessplating on spreading kinetics and contact angles of flatcoated-ceramics using a sessile drop technique [32,33], per-mitted to establish a temperature range of 800–900 �C toinvestigate the influence of Ni-coating on the infiltrationbehavior of the compacts. Generally speaking, it isreported that nickel and copper improved wettability andadhesion with aluminum resulting from the dissolution ofthe coating in the liquid drop and the exothermic forma-tion of intermetallic compounds.

3. Metallization of SiC powders

The silicon carbide coated-powders were inherently het-erogeneous, in that each particle consists of a core of cera-mic material and a metallic coating. The composition ofthe bulk Ni-coated powder deposits determined by ICP-

AES analysis ranged from 1.6 to 1.9 wt% for the phospho-rus content, 17–21 wt% of metallic nickel, and the balancewas ceramic. The theoretical density of the coating andthus, the density of the bulk Ni-coated composite powderswere evaluated through the rule of mixtures. From thechemical composition, the calculated density for the Ni-film ranged from 7.8 to 8.0 g/cm3, while the bulk densityof the Ni-coated SiC powders was 3.6 g/cm3.

The structure of the nickel coating on the ceramic sur-face consisted of fine-grain deposits coalescing continu-ously to form the metallic film. A continuous nickel filmwas observed over the whole surface regardless of sizeand shape. XRD patterns of the Ni-coated powdersshowed a pronounced broad band in the region corre-sponding to the Ni (111) and (200) reflections, signifyingthat the Ni–P deposits contained the metal in an amor-phous form. The amorphous structure of the nickel filmsis due to the presence of phosphorus in the deposit.

4. Composite fabrication and characterization

Infiltration results confirmed the favorable effect of sur-face modification during composite processing. The metal-lic coating improved the wettability of ceramics, promotingeasy metal flow through the preform in comparison to thenon-infiltration behavior of the uncoated counterpart sam-ples. At the same time, it changed the matrix compositionby alloying with the matrix. Because the preforms were notpreheated prior to immersion in the aluminum bath, it tooksome time for them to achieve equilibrium with the temper-ature of the melt. The duration of such an incubationperiod was in the order of tens of seconds, and decreasesconsiderably at higher infiltration temperatures. Infiltra-tion of the compacts was attempted from temperatures aslow as 700 �C. At this temperature, the viscosity of thematrix was high compared to higher temperatures. Henceinitial liquid infiltration was slow. This incubation periodwas long and significant scatter in the data was observedbut the samples were eventually infiltrated. It is possiblethat the infiltration kinetics slowed due to the solidificationof molten metal when it came in contact with the cold Ni–SiC particulates. This led to the formation of a thin solid-ified layer on the particulates, thus reducing the overallpore volume available for infiltration.

The infiltration length achieved as a function of timeand temperature for Al/Ni–SiC composites is shown inFig. 1 for Ni-coated 78 lm SiC reinforcements. The infil-tration distances plotted are averages of several sectionedsamples measured. A marginal increase in the infiltratedlength was evident with an increase in melt temperaturewhich could be attributed to the decrease in the surface ten-sion of the Al melt and the favored reactions between Aland Ni, which resulted in a decrease in the contact anglebetween the particulates and Al causing improved wettabil-ity. Infiltration of the preform was aided by the exothermicreaction occurring upon formation of nickel aluminide,NiAl3 (�DH = 150.7 kJ/mol). The dissolution of the

0

10

20

30

40

50

60

70

0 30 60 90 120 150

Time (sec)

Len

gth

(m

m)

800 °C 850 °C 900 °C

Fig. 1. Infiltration length of aluminum into coated Ni–SiC.

214 C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218

coating, creating locally clean SiC surfaces at the bottom ofthe compacts, along with the partial saturation of Al limit-ing the reaction between Ni and Al at the top of the com-posites, would slow down melt infiltration. Thus, theinfiltration curves followed a classical parabolic trendirrespective of particle size and temperature. Particulatecomposites reinforced with nickel-coated SiC were infil-trated at 900 �C in less than 2 min, while in the case ofthe as-received powders, no infiltration was observed after20 min under similar experimental conditions.

Liquid metals almost always wet solid metals and a largeimprovement in wetting occurs by formation of intermetal-lics compounds (e.g., Al–Ni). By X-ray diffraction, thepresence of NiAl3 was detected in Ni-coated SiC (Fig. 2).

20 25 30 35 40 45

2θ (deg

Inte

nsi

ty (

arb

. un

its)

Fig. 2. XRD trace of Al/Ni–SiC com

Other phases were present in small proportion such asNiAl or NiAl2, which were confirmed by quantitativemicroanalysis. The results of X-ray diffraction at differentpositions along the infiltrated samples revealed no evidenceof nickel in the matrix at the very bottom of the samples.This leads to the conclusion that the nickel was dissolvedaway from this section of the composites. The aluminumstream thus enriched with nickel solidifies near the top ofthe sample, leading to the precipitation of the nickel richphases. This creates a concentration gradient of nickelfrom the bottom to the top of the infiltrated sample. Suchconcentration gradients have been reported when infiltrat-ing coated-reinforcements with molten aluminum [20,22].Regarding the distribution of reinforcement phase in com-posites fabricated from Ni-coated particulates, Chung et al.[21] reported SiC free zones close to the edge of the speci-mens when reinforcing aluminum with 20 vol.% Ni coatedSiC particles mixed with pure Al powders. In the presentcomposites no SiC free zones were observed. The micro-scopic distribution of SiC particulates was quite homoge-neous along the infiltration axis in all the regions of theinfiltrated samples.

Fig. 3 shows microstructural features of the Al-rich andNi-rich matrix zones of the composites. In some areasaround the particles in the top section, a double layer inter-facial zone was observed with entrapped matrix pockets(Fig. 3b). Evidence of interfacial debonding due to solidifi-cation shrinkage is reduced since it is reported that theliquid to solid shrinkage in Al–Ni alloys is also reduced lin-early so that at 15% Ni, it drops to about 1% [34]. Noapparent interfacial reaction between the matrix and theceramics was detected. Thus, the formation of the detri-mental Al4C3 was avoided in the composites reinforcedwith silicon carbide particles. However, the presence ofsmall black precipitates was observed in the matrix in some

50 55 60 65 70

rees)

SiC Al NiAl3 Si AlP

posite along the infiltration axis.

Fig. 3. Microstructures observed in (a) Al-rich and (b) Ni-rich matrix zones.

C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218 215

regions of the composites. Energy Dispersive X-ray spec-troscopy (EDS) of these phases showed that they containphosphorous. In addition to P, O was also present in somephases. In general, the Al–P precipitates were clusteredaround the SiC particles.

By measuring the density of sectioned composites, a gra-dient was found along the infiltration axis of the samples.Fig. 4 shows the increase in density along the length ofthe Al/Ni–SiC composites when infiltrated at 800 �C. Therise in density at the top of the composites was due tothe segregation of nickel from the stream of aluminum dur-ing infiltration, leading to precipitation of Ni rich phases

(the density of NiAl3 and Ni2Al3 are 3.95 and 5.15 g/cm3,respectively, compared to 2.7 g/cm3 for pure Al). Inaddition, the samples reinforced with 78 lm Ni–SiCparticles exhibited a higher density in all sections of thecomposites than the 49 lm Ni–SiC reinforced materials.This is attributed to the higher ceramic content in the78 lm SiC composites, 45.7 vol.%, compared to the lowerSiC content, 42.8 vol.%, in the fine particulate composites.Rohatgi et al. [35] previously reported density gradientsin MMCs due to solute segregation during the infiltra-tion of liquid aluminum in beds of nickel-coated fly ashcenospheres.

2

2.4

2.8

3.2

3.6

4

0 10 20 30 40 50 60

Position from bottom, mm

Den

sity

, g/c

m3

78 μm

49 μm

Fig. 4. Density at different positions along the infiltrated Al/Ni–SiCcomposites.

Fig. 5. Typical fracture modes observed in Al/Ni–SiC composites. (a) A brittle

216 C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218

Similar to density, hardness of the composites increasedfrom bottom to top along the infiltration axis. The bottomportion of the composites showed a hardness value in theorder of 100 HV, corresponding to a nickel free matrix.Generally, the hardness in the center of the samples wasmaintained and nearly unvarying, increasing near the topbecause of the presence of intermetallic phases and Al–Nieutectic islands. The surface of the composites showed aclear rise in hardness upon approaching the top of thebar, having an average value in the range 250–300 HV.

The presence of �45% particles in the Al/Ni–SiC com-posites led to a substantial increase in the Young’s modulusof the matrix. The Young’s modulus of the 78 lm SiC com-posites was close to 160 GPa, while in the case of the 49 lmSiC composites it ranged from 145 to 150 GPa (comparedwith 70 GPa for the unreinforced aluminum). Flexurestrengths of composites reinforced with 78 lm SiC werein the range of 165–205 MPa, whereas the materials rein-

-ductile fracture region (800 �C) and (b) a brittle fracture region (800 �C).

C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218 217

forced with 49 lm SiC particles had higher values in theorder of 200–225 MPa, i.e., a nearly 15% improvement ofstrength. These results indicate a large scatter with infil-tration temperature for the large particle reinforced materi-als. However, the mean bend strength values are higherthan those reported by Chung et al. [21,22] for Al/Ni–SiC composites, who found strengths in the order of 120–150 MPa using Al–Si alloys as the matrix. The higherMOR value of the small particle reinforced composite isattributed to the smaller particulates which lead to sur-faces free from large defects compared to the large rein-forcements.

Observations of the fracture surfaces via SEM show thecomposites to exhibit a combination of completely brittlefracture regions and brittle-ductile fracture regions. Frac-ture surfaces were flat and normal to the stress axis whenviewed on a macroscopic scale but rough when viewed ona microscopic scale. In the brittle-ductile fracture regions,microdimples associated with Al enriched zones wereobserved in the matrix (Fig. 5a). The presence of dimplesindicates that ductile failure had occurred. As shown inFig. 5b, no matrix microdimples were found in the com-pletely brittle fracture regions and cleavage occurred inboth the matrix and particulate reinforcement. The frac-ture surface with fewer dimples was associated more withparticle cleavage and less pull out. Thus, the brittle failuremode corresponds to the fracture of SiC and intermetallicsrich zones; NiAl3 leads to a reduction in material ductility,particularly given its high content and gross morphology.A shear punch test was used as an alternative for assessingthe tensile properties of the Al/Ni–SiC composites [36].Fig. 6 shows the tensile properties at room temperatureof composites fabricated at 800 �C with either 49 or78 lm SiC particles. It is observed that upon approachingthe top of the infiltrated bars, the strength of the compos-ites increased in response to the higher content of Al–Ni

150

200

250

300

0 10 20 30 40 50 60

Distance from bottom, mm

UT

S (

MP

a)

78 μm

49 μm

Fig. 6. Variation of the tensile strength along the length of Al/Ni–SiCcomposite bars from shear punch test determinations.

phases. Since very small elongations to fracture wereobserved, the UTS values were found to be similar for allmaterials. The strength of the thin specimens went from�220 MPa for samples with a low Ni content matrix tonearly 250 MPa for the enriched Ni matrix portions ofthe composites. The tensile properties calculated arerelated to both, the high volume fraction of the hardceramic phase in the samples, and the formation of NiAl3intermetallic. Considering that the tensile strength of theunreinforced Al matrix is 45 MPa, an increase in strengthof 500% was obtained for the infiltrated composites.Such an increase is substantial, since increases of only60% in yield and ultimate tensile strengths, depending onthe volume fraction of reinforcement, the type of alloy,and the matrix alloy temper, are normally reported forthe strength of discontinuously reinforced Al alloys [28].However, the strength of the composites was lower thanthat reported for similar materials produced by conven-tional systems due to the coarser particle size used. It isexpected that with some microstructural and processrefinement, the present technique could lead to higher val-ues of strength.

5. Conclusions

Electroless plating can be considered as an intermediatestep in the fabrication of ceramic–metal composites, sinceit enhances the wettability of the ceramic phase for com-posite fabrication. The use of metal-coated reinforce-ments promoted easy metal flow through the preforms,compared to the non-infiltration behavior of the uncoatedcounterpart samples. Reduction of the SiC/Al interactionand enhancement of wetting can be achieved by modify-ing the ceramic surface. Infiltrated bars of 6 cm lengthand 8 mm diameter were infiltrated in times as short as2 min. XRD microstructural analysis of the compositesindicated the formation of intermetallic phases such asNiAl and NiAl3 when nickel-coated powders were infil-trated. The density, hardness and strength of the compos-ites increased as the overall nickel concentration of thematrix increased from bottom to the top of the samples.By controlling the initial content of metal on reinforce-ments and the solidification conditions, the structure andproperties of composites can be tailored to fit differentengineering applications.

Several issues remain with regard to the use of interlay-ers in composite processing. Of major significance is con-trolling the amount and composition of the coatings.Recently it has been shown that other plating bath chemis-tries may be used to enhance the performance of the resul-tant composite microstructure. Moreover, the applicationof electroless coatings to control hydride decompositionduring metal foam production has been demonstratedand such coatings have been applied in various advancedjoining processes. Therefore the potential scope for thistechnology is very promising in a wide variety of techno-logical fields.

218 C.A. Leon-Patino, R.A.L. Drew / Current Opinion in Solid State and Materials Science 9 (2005) 211–218

Acknowledgements

The authors would like to thank to McGill University,CONACYT-Mexico and Universidad Michoacana, Mex-ico (UMSNH), for the financial support of this research.

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