Ni–Al–Ti coatings obtained by microwave assisted SHS: Effect of annealing on microstructural and...

Surface & Coatings Technology 203 (2009) 1429–1437

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Ni–Al–Ti coatings obtained by microwave assisted SHS: Effect of annealing onmicrostructural and mechanical properties

G.P. Cammarota a, A. Casagrande a,⁎, G. Poli b, P. Veronesi b

a SMETEC-University of Bologna, Italyb Department of Materials and Environmental Engineering, University of Modena and Reggio Emilia, Italy

⁎ Corresponding author.E-mail address: [email protected] (A. Casagra

0257-8972/$ – see front matter © 2008 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2008.11.017

a b s t r a c t
a r t i c l e i n f o
Article history:
Microwaves at 2.45 GHz hav Received 15 April 2008Accepted in revised form 19 November 2008Available online 3 December 2008
Keywords:CoatingsMicrowavesInterlocking microstructureLiquid phase bonding

e been used to ignite the self-propagating high-temperature synthesis (SHS) in aNi and Al powder mixture to produce NiAl coatings on Ti. The temperature of the reacting powders, owing tothe exothermic nature of the reaction and to the energy provided by the high frequency electromagneticfield, reaches the melting temperature of NiAl. The newly formed NiAl, in the liquid phase, reacts with theunderlying Ti and rapidly forms at the interface a ternary eutectic belonging to the Ni–Al–Ti system. Duringcooling, the formed transient liquid eutectic layer at the Ti/NiAl interface originates an intermediate layerpresenting an interlocking network, which constitutes an in-situ reinforcement phase. In this study it isshown that the thickness of the interface layer can be altered by annealing at temperature higher than theternary eutectic temperature (904 °C), in a few minutes time. Experimental results demonstrated also that,due to its peculiar microstructure, the intermediate layer is tougher and harder than the parent NiAl coating.

© 2008 Elsevier B.V. All rights reserved.

1. Introduction

Titanium and its alloys are suitable candidates for fabrication oflight weight components to be used in the aerospace and automotivesector, owing to their high-temperature strength in combination withlow density [1]. However, due to severe oxidation problems at hightemperature (TiO2 forms at temperature above 500–520 °C) andreduced wear resistance, technical applications of Ti and Ti-alloys arecurrently limited to temperaturesmuch lower than those theoreticallyallowed by their melting point and mechanical properties [2]. As amatter of fact, the service temperature limit set by the mechanicalcapabilities of most advanced titanium alloys is 520–600 °C. Improve-ment of the high temperature oxidation and wear resistance oftitanium alloys can be achieved by a modification of the surfaceproperties, for instance applying protective coatings which helpmaintaining the excellent bulk properties of the substrate [3]. Duringlast decades, many studies have been performed on surface treat-ments to improve oxidation, wear and corrosion resistance: overlaycoating [4,5], Al-Ti and Al–Ni intermetallic phase obtained by packcementation [6] and hot-dipping diffusion coating [7], silicide/ceramics coatings [8], plasma spraying techniques, electron beamphysical vapour deposition of intermetallic compounds [9,10], ionimplantation [11,12], laser surface modification [13] and surfaceoxidation [14]. However, many of these coatings are still far fromsatisfactory [15], especially as far as adhesion is concerned. The

nde).

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importance of achieving a strong and resistant bond to the substrate isthe first requisite in producing durable coatings.

In many cases, substrate and coating material combinationsrequire the presence of a supplementary interface-layer to promotebonding. With this respect, powder metallurgy processing, and inparticular the Self Propagating High Temperature Synthesis (SHS), hasproved to be an interesting route to form protective coatings ontitanium [16]. The great exothermicity of this synthesis route and thehigh temperature reached promotes the existence, for a limited periodof time, of a highly reactive molten phase [17].

In this work, a novel two step technique to develop thick protectivecoatings on titanium, having a strong adhesion to the substrate, ispresented.

In the first step, an outer NiAl coating and an intermediate Ni–Al–Ti eutectic interface layer are formed on Ti substrate, using micro-waves at 2.45 GHz to ignite and sustain the SHS of Ni and Al powdermixture [16]. In the second step, the thickness of the interface layer isprogressively increased by annealing at temperature higher than theternary eutectic temperature (904 °C), for times in the 5–20min range.

In the first step, the temperature of the reacting powders, owing tothe exothermic nature of the reaction, reaches the melting point ofNiAl [18], which then rapidly forms a ternary eutectic with theunderlying titanium. The newly formed eutectic remains in the liquidstate due to the high temperatures developed by the powdersreaction, and to the energy still supplied by the microwave field.Thus, a reactive liquid interface between the titanium substrate andthe NiAl coating is maintained even after synthesis. As a matter of fact,numerical simulation has shown that microwaves are able to convey

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Fig. 1. Optical micrograph of the cross section of the as synthesised NiAl coating (right)on Ti substrate (left), with the presence of the intermediate ternary layer (centre).

1430 G.P. Cammarota et al. / Surface & Coatings Technology 203 (2009) 1429–1437

energy to the reacting system during and also after SHS ignition, thusextending the time of existence of the liquid reactive interface [18].This improves adhesion and promotes the formation of a strongchemical bonding [19]. During cooling, the formed transient liquideutectic layer at the Ti/NiAl interface originates an intermediate layer,presenting an interlocking network which constitutes an in-situreinforcement phase. The interlocking structure and the fine grain sizehave been widely studied in weld metals [20–24] and heat affectedzones; such microstructure resulted effective in increasing bothstrength, toughness and resistance to fracture propagation [25,26].

The aim of the present work is to study the mechanical propertiesand chemical bonding of the interlayer formed between the NiAlintermetallic coating and the Ti metallic substrate. In particular, theeffect of micro-mechanical interlocking with non planar contacts atthe Ti–Ni aluminide coating interface is investigated, as well as itseffects in improving adhesion and strength of the coating. Thedistribution of elements across the interface was studied as well, inorder to determine the degree of homogeneity of the intermediatelayer as a function of the annealing time at 950 °C, and subsequentcooling rate.

2. Experimental

Ni powders (99% pure, 5–7 µm in diameter) and Al gas atomizedpowders (99.8% pure, 10–50 µm in diameter) in 1:1 atomic ratio weremixed in an alumina jar for 30 min under vacuum. Successively, 1 g ofthe mixed powders was placed on a CP titanium disc(diameter=12 mm, thickness=5 mm) and uniaxially pressed at140 MPa.

Thegreen couple (titanium+compactedpowders)was inserted intoa microwave single mode applicator, to start the microwave assistedSHS. In order to reduce the porosity after SHS, in some configurations asandwich structurewas prepared, adding a further titaniumdisc on thetop of the compacted powders. This arrangement allowed to apply amoderate load (0.01 MPa) during synthesis, using a quartz rod incontact with the upper disc. The use of quartz rod in the microwaveapplicator allowed to avoid any significant perturbation of theelectromagnetic field distribution. A microwave TE10n single modeapplicator (Transverse Electricmodewhere the indexes describe a fieldconfiguration with 1 semi-period variation of the electric field alongthe applicator width, 0 variations along the applicator height, and nvariations along the applicator length), based on the WR340 rectan-gular waveguide geometry (86×43mm section), operating at 2.45 GHzwas used to start the SHS. The sample was placed, one per run, on analumina refractory plate, in order to avoid contact with the applicator'swalls and to provide thermal insulation. The heat is generated only bymicrowave interaction with the compacted powders, while themetallic substrate does not couple with microwaves. A 20 ml/min Arflux prevented the oxidation of the powders. In case of load matchedand using a 1200 W emitted power, the microwave assisted SHSstarting and completion required a maximum time of 30 s. Power wasturned off one second after SHS ignition, in order to produce anintermediate layer nomore than 50 µm thick, according to themethoddeveloped in a previous study [27]. The coated disc was removed fromthe applicator after 300 s from the start of the exothermic reaction.Each sample was divided into four parts and each one was annealed at950 °C in a resistance furnace under argon flux, respectively for 5,10,15and 20 min. Two different cooling schedules have been applied afterannealing: slow cooling in the furnace at 5 K/min and rapid cooling instatic air at approximately 50 K/min.

Microstructural evolution of the coating/Ti intermediate layerthickness was investigated by optical microscopy after glyceregiaetching and by scanning electron microscope ZEISS EVO 50. Phaseanalysis was carried out by a layer by layer X-ray diffraction technique(PHILIPS PW 1720) using CuKα radiation. Elemental chemical analysiswas carried out using EDS (Energy Dispersive X-ray microanalyzer

Spectroscopy) OXFORD INCA ENERGY 350. DSC analysis (PL-STARHEOMETRIC-SCIENTIFIC) was used to determine the ternary eutectictemperature. Vickers micro- and macro-indentation (ISOSCAN GALI-LEO) was used to measure samples (micro)hardness and to determinetheir fracture toughness according to Evans Equation [28].

3. Results and discussion

3.1. Microstructure of coatings obtained by microwave-assisted SHS

When exposed to microwaves at 2.45 GHz, the Ti and Ni–Alpowder compact react and the SHS reaction starts to formintermetallic compounds. During the exothermic reaction, heat istransferred from the newly formed melted NiAl to the Ti surface.Due to the extremely rapid reaction and the relatively low thermalconductivity of titanium, the extension of the heat affected zone issmaller than 80 µm, as demonstrated in a previous study [18]. Theuse of microwaves allows to continue to provide energy to thenewly formed phases, despite the adverse temperature gradient(the synthesised compounds are at temperature much higher thanthe surrounding environment). This reduces the cooling rate andallows to extend the time of existence of the molten phase [29].

These peculiar heating and cooling conditions promote theformation of a duplex microstructure (Fig. 1) constituted of [17]:

1) a polycrystalline outer layer essentially made of NiAl grains,without preferential orientation

2) an intermediate Ni–Al–Ti layer that in the applied synthesisconditions, has a 50 µm average thickness and presents a dendriticand eutectic habitus

The nature of the intermediate layer can be explained asfollows. Notely, in the binary high Ti content side of the Ti–Nisystem, the final microstructure is due to a eutectoidic transforma-tion of a Tiβ–Ni solid solution into Ti2Ni plus α-Ti (left side of Fig. 2[30]). The eutectoidic transformation occurs in the solid state, at765 °C, which is a much lower temperature than the ternaryeutectic temperature in the Ti–Ni–Al system. Due to the highaluminium content of the starting powders, in the intermediatelayer is expected the presence of nonequilibrium Ti2Ni and α-Tiphases presenting a variable Al content.

The appearance of the liquid phase in the intermediate Ti–Ni–Allayer at temperatures above 904 °C, as determined by DSC analysis(Fig. 3), allowed to design subsequent post-coating annealingtreatments with the purpose of altering the intermediate layerthickness. As a matter of fact, the presence of a liquid phase isexpected to promote further reaction with the NiAl and Ti

Fig. 2. Ti–Ni phase diagram [30].

1431G.P. Cammarota et al. / Surface & Coatings Technology 203 (2009) 1429–1437

surrounding the intermediate ternary layer. Annealing treatmentswere thus conducted in air at 950 °C for 5, 10, 15 and 20 min, inorder to increase and control the intermediate layer thickness,varying the subsequent cooling conditions to try to preserve itsmicrostructure.

Fig. 3. Differential scanning calorimetry (DSC, 10 °C/min heating rate) of intermedia

3.2. Microstructure of the coatings after heat treatments at 950 °C

Annealing treatments at 950 °C allowed to progressively increasethe intermediate layer thickness, as shown in Fig. 4. The growth of theintermediate layer occurred at the cost of both NiAl and Ti, whose

te ternary eutectic layer to evaluate the liquid phase appearance temperature.

Fig. 4.Microstructure of the intermediate layer after annealing at 950 °C for: a) 5 min; b) 10 min; c) 15min; d) 20min. Low cooling rate was applied after annealing. From left to right:NiAl, intermediate layer, Ti substrate.


thickness is reduced, indicating that the solid–liquid interface movestowards both the solid NiAl coating and Ti substrate.

Fig. 4a) shows the interface layer microstructure after 5 min ofannealing at 950 °C, indicating that there has been a pronouncedincrease of the intermediate layer thickness, which reached almost200 µm, starting from the initial 50 µm. On the right of eachmicrograph of Fig. 4 there is the Ti substrate, while on the left there isthe newly formed NiAl coating. Similarly, in Fig. 4b) a further thicknessincrease of the intermediate layer is evident after 10 min annealing,now reaching almost 250 µm, which goes up to 300 µm after 15 min(Fig 4c). and in excess of 400 µm after 20 min. Almost identicalmicrostructures are obtained after the different times of annealing at950 °C and low cooling rate (5 K/min), with the only difference that,after 20 min, the morphology changed from a lamellar eutectoidicstructure near Ti interface to spherical coarsened particles near theNiAl interface. These microstructures are promoted by the aforemen-tioned eutectoidic transformation in the Ni–Ti system, favoured by thelow cooling rates applied.

From the annealing treatments, it resulted that in order to controlthe intermediate layer thickness and microstructure, the ternaryeutectic temperature and the annealing holding time are the mostimportant parameters. However, cooling rate plays an important role

Fig. 5. Optical micrograph of the cross section of the intermediate ternary layer after20 min annealing at 950 °C followed by rapid cooling. From left to right: NiAl,intermediate layer, Ti substrate.

too, as shown in Fig. 5. In case of high (50 K/min) cooling rate afterannealing at 950 °C for 20 min, a different microstructural morphol-ogy is obtained with respect to the slow cooling rate one (Fig 4d),presenting tiny interlocking dendritic shapes. High cooling rates tendto inhibit the eutectoidic transformation in the Ti–Ni system and theresulting microstructure consists typically of small dendrites, inter-dendritic eutectoid and dispersed α-Ti (Ni, Al) particles. Despite the

Fig. 6. a) SEMmicrograph (BSE): cross-section showing the eutectoid andα-Ti phase; b)optical microscope micrograph of cross-section showing the eutectoid and α-Ti phase.

Table 1Composition of the Ti/NiAl interface layer after SHS reaction and subsequent heat-treatments at 950 °C

SEM Micrographs of the cross sections and semiquantitative EDS analysis.


Table 2XRD layer by layer spectra of the eutectoidic intermediate ternary layer after annealing at 950 °C for 20min starting fromα-Ti interface (top) to NiAl interface (bottom); a) low coolingrate; b) high cooling rate


Fig. 7. SEM micrograph (BSE) of a cross section at the NiAl/intermediate layer interface,after annealing at 950 °C followed by low cooling rate, showing crack not propagatingfrom NiAl (left) to the ternary layer (right).

Table 3(Micro)hardness and indentation fracture toughness of the NiAl coating and intermediate layer on titanium, as a function of annealing time and subsequent cooling rate (c.r.)conditions

Intermed.layer asSHSproduced

Intermed. layer H.T. 950 °C For5 min





NiAl Coating H.Treat. at 950 °C20 min

Ti sub. Interf. H.T.at 950 °C 20 min.

Low c.r. Low c.r Low c.r Low c.r High c.r Low c.r. Low c.r.

Micro-hardnessHV0.05

1020±70 967±95 955±120 982±105 934±140 1072±84 505±26 480±16

Hardness HV 722±48 702±83 689±91 711±86 714±112 738±50 298±22 277±13Identation fracturetoughness MPa m1/2

3.5±1 4.4±1 1.2±0.1


different microstructure, the intermediate layer thickness is almostidentical to the one obtained by slow cooling, i.e. slightly more than400 µm. Figs. 4 and 5 show also that there is good metallurgicalbonding at both sides of the intermediate layer, which could beascribed to the crystalline structure of the constituents. Thecontinuous crystallographic structure (α-Ti is found in the substrateand also in the intermediate layer after cooling), is likely to improvethe bonding strength, whichmakes the coatingmore prone to toleratethermal stresses normally induced by the different thermal expansioncoefficients of its constituents. As a matter of fact, a crystallographiccontinuity tends to improve adhesion with respect to the case of twodifferent crystalline structures in strict contact.

The tiny dimension of dendrites and lamellae, shown in Fig. 6, gives anindication of the small diffusion distance involved during cooling in thisintermediate layer.

It is evident a varying lamellae spacing, probably because of thedifferent orientation of eutectoidic colonies. It should be noted thatthe size and the volume fraction of dendrites, and/or of lamellaeeutectoidic structure depend on the cooling rate experienced after theisothermal treatment and both of these factors are expected to affectthe strength, the hardness and the toughness of the intermediatelayer.

3.3. Phase identification

The chemical composition of constituents in the intermediate layerbetween NiAl and Ti substrate, was determined by semi-quantitativeEDS analysis as a function of annealing times and subsequent lowcooling rate. Results (average of 6 measures in similar regions),together with SEMmicrographs (BSE) indicating the analyzed regions,are summarised in Table 1.

The monolithic coarse dendritic phase in Table 1 contains largeamounts of dissolved Al in α-Ti matrix and smaller amounts of Ni, as aconsequence of partial Ti dissolution in contact with eutectic liquidphase. This is in agreement with phase diagram of Ti–Ni and Ti–Alsystems [30].

EDS semi-quantitativeanalyses carriedouton the intermediate layerofthe 20 min annealed samples subjected to rapid cooling, show only anaverage chemical composition of the dendritic-interdendritic eutectoidicstructure similar to that obtained in case of low cooling rate. Due to thesmallness of the structures, from a morphological point of view it isimpossible todistinguishbetweeneutectoidic structureandα-Tiphase (orß-Ti phase); thus, identification of the phases constituting the dendriticintermediate layer required the use of X-ray diffraction.

XRD layer-by-layer data analysis, reported in Table 2, referred to20 min annealed samples, subjected to low and high cooling rate,confirmed that hexagonal α-Ti (Al,Ni) and cubic Ti2Ni (Al) composethe intermediate layer. No other secondary phases, nor unreactedconstituent elements are detected. The comparison of the X-raydiffraction results with the elemental composition values obtained byEDS analysis, confirmed that the eutectoidic constituents and the α-Tisolid solution are contained also in the intermediate layer annealed at

950 °C and cooled at high cooling rate. The diffraction patterns of bothseries of spectra (high and low cooling after annealing at 950 °C) areidentical with same trends: high volume fraction of α-Ti (Al,Ni) phaseis present near Ti interface and high volume fraction of Ti2Ni (Al)compound is found near the NiAl coating interface, depending on theNi at.% composition (as deducible by the Ni–Ti system).

In agreement with Huneau et al. [31], the multi-phase micro-structure corresponds to an hypereutectoidic composition rangingfrom Al15Ti56Ni29 (bright areas of figures inserted in Table 1) to finelamellar hypoeutectoidic area Al24Ti58Ni18 (grey areas of figuresinserted in Table 1). As reported by J.C. Schuster [32], Ti2Ni phasedissolves, in equilibrium conditions, up to 14 at.% Al at 900 °C, whereAl substitutes for Ti.

Our EDS data present small differences of the phase compositions,probably due to the fact that solidification takes place in nonequilibrium conditions, starting from temperatures well above900 °C. The origin of the phases composition observed in the presentwork, in the Ti-rich range composition, can be ascribed to the fact thatthe eutectoid reaction β-Ti=α-Ti (Al,Ni)+Ti2Ni (Al) occurs, duringcooling, at temperatures lower than 765 °C, due to Ni influence [33].Moreover, α-Ti phase formation is favoured by the abundance of Al,which is a well known α-stabiliser, accompanied by elevatedelemental diffusion phenomena, typical of binary Ni–Ti alloys [30].On the other hand, XRD spectra layer by layer, carried out on theintermediate layers after the annealing at 950 °C at high cooling rate,do not attest the presence of the ß-Ti phase, notwithstanding thesignificant Ni at.% (as ß-stabilizer), and the possibility that thedendritic nucleation of the ß-Ti phase, at high cooling rate, is favouredwith respect to an exagonal α-Ti phase. These considerations are in


favour of the existence of a ß-Ti phase or of transformation products ofthe ß-Ti phase at high temperatures. For this reason, other analyses,such as TEM analysis, should be necessary to better investigate thestructure of the intermediate fine dendritic intermediate layer.

3.4. Relationship between eutectoidic microstructure and mechanicalproperties of the intermediate eutectoidic layer

The microhardness HV0.05 and hardness HV values of the aluminidecoatings before and after annealingweremeasured along the transversesection of the polished samples and are reported in Table 3. The resultsshow that the intermediate layer with hypoeutectoidic structure has anaverage microhardness ranging from about 790 HV0.05 near the Tiinterface (where thehighest presence ofα-Ti phase is expected to occur),to about 1150 HV0.05, hypereutectoidic structure, near the NiAl interface(where the prevailing microstructure is based on the co-existence ofeutectoidic constituents and Ti2Ni). At low and high cooling rates, thesamples present similar trends in (micro)hardness values across thetransverse section of the interface layer; generally higher for the finedendritic microstructure morphologies (obtained with high coolingrate).

In all cases, the intermediate layer presents (micro)hardness valueshigher than the NiAl coating (505±26 HV0.05 and 298±22 HV). As amatter of fact, the network of interlocking fine dendritic structuresand the small distance in the eutectoidic lamellar structure play animportant role in the hardness improvement.

Indentation fracture toughness of NiAl coating and intermediateternary layers with eutectoidic structure was measured by the Evansand Charles Method [28]. From the fracture toughness measurements,two trends are evident:

• a higher toughness was achieved in the ternary layer witheutectoidic structure: average fracture toughness value of ternaryintermediate layer is 4.4 and 3.5 MPa m1/2 for high and low coolingrate respectively, compared to 1.2 MPa m1/2 of the NiAl coating. Theobtained results confirm the tougher nature of the intermediatelayer evidenced also in the optical microscopemicrographs of Figs. 7and 8.

• the intermediate layer indentation toughness varies widely in theanalysed specimens and this can be ascribed to the local inhomo-geneity of the microstructure, which includes hard constituents

Fig. 8. Optical microscope micrograph of a cross section at the NiAl/intermediate layerinterface after annealing at 950 °C followed by high cooling rate, showing crack notpropagating from NiAl (left) to the ternary layer (right).

(eutectoid) toughened through the incorporation of more ductilephases which stretch, neck and plastically deform in the presence ofa propagating crack tip. This behaviour is typical of the CrackBridging mechanism [34].

It has been generally found that indentation toughness measure-ments do not agreewell withmeasurementsmade using conventionalfracture mechanics specimens [35]. Some of the discrepancies can berelated to the precision of the expressions, to the failures in specimensand to the effects such as surface residual stresses or various crack/microstructure interactions. However, the obtained results show goodagreement with the NiAl toughness from literature data [35].

The measured increase of hardness and fracture toughness can beexplained considering the intermediate layer composition andmicrostructure. As shown in the previous paragraph, at low coolingrate after annealing, the general microstructure of the intermediatelayer is prevalently composed by a ternary eutectoidic structure withsmall lamellae dimension, together with a third plastic phase: α-Ti(Ni,Al). Increasing the cooling rate allows the formation of a finerdendritic interlocking microstructure. Thus, in the studied system, thestrength and the toughness mainly come from:

a) eutectoidic matrix, α-Ti (Al,Ni)–Ti2Ni, which hinders crackspropagation and presence of ductile phase (α-Ti) which providesthe capacity for large plastic strain, thus improving the resistanceto fracture propagation (Fig. 7); moreover, the presence of multipleinterfaces between such different constituents promotes crackdeflection. The minimization of the energy required for the crackto propagate through the network of different structures is thenachieved by changing the propagation direction and avoiding thetougher regions.

b) anisotropic dendrite growth orientations in the interlocking finedendrite microstructures which can lock the crack propagation(Fig. 8)

The improvement of the fracture toughness with respect to thebrittle NiAl coating can be also explained on the basis of the effects ofresidual stresses due to thermal expansion anisotropy, differencesbetween each constituent thermal expansion coefficients, non planarcontacts between anisotropic dendrite secondary arms and thepresence of a plastic dendritic or dispersed α-Ti(Ni,Al) phase.Moreover, there is experimental evidence that, in addition to residualstress effects of the type considered in this work, crack bridging byuncracked dendrites and fine anisotropic grain size growth anddendrite-interlocking effects, formed during the eutectic isothermalsolidification in a short time, also serve as toughening mechanism inpolyphase interlayer at the metal-coating interface [33].

The obtained interlocking and eutectoidic microstructure, with theabove mentioned mechanical properties, on the basis of the obtainedresults, might replace the functionality of hard and brittle NiAlaluminide coatings on metallic Ti substrate in service, increasing thewear resistance and the toughness [36]. However, further investiga-tion should be conducted regarding the capability of the ternary Ni–Al–Ti layers to replace the intermetallic NiAl coating at hightemperatures, up to 900 °C.

4. Conclusions

Intermetallic NiAl-based thick coatings, resulting frommicrowave-assisted SHS of elemental Al and Ni powders on Ti substrate wereobtained. The capability of microwaves to convey energy to thereacting species, during and after SHS, favoured the formation of anintermediate ternary layer between the newly formed polycrystallineNiAl and the Ti substrate. This process provides a simple robustbonding solution with rapid processing time in the production of aduplex microstructure coating on titanium substrate. The intermedi-ate layer belonging to the Ni–Al–Ti system exhibits a dendritic–


eutectoidic morphology. In the applied synthesis conditions, theintermediate layer had a thickness of 50 µm. Subsequent annealing at950 °C for 5 to 20 min allowed to increase the intermediate layerthickness up to 400 µm, preserving its interlocking microstructure.Intermediate layer growth occurs at the cost of both Ti substrate andNiAl outer coating.

It was found that hardness and fracture toughness of theintermediate layer is much higher than the ones of each singleconstituent and of the NiAl coating, due to the strong constrainingeffects of the interlocking dendrite-lamellar microstructure.

The obtained thick intermediate layer could replace the function-ality of hard and brittle NiAl aluminide coatings in service withimproved mechanical properties, like hardness and fracturetoughness.

Acknowledgements

The authors wish to thank Dr. F. Tarterini and Ing. F. Corvucci of theSMETEC Department of Bologna-University for helpful discussionsand accessibility to EDS-SEM and XRD equipment.

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