Thermal stability of nanoscale metallic multilayers

    Thermal stability of nanoscale metallic multilayers A.S. Ramos, A.J. Cavaleiro, M.T. Vieira, J. Morgiel, G. Safran PII: DOI: Reference:

S0040-6090(14)00640-3 doi: 10.1016/j.tsf.2014.05.065 TSF 33506

To appear in:

Thin Solid Films

Received date: Revised date: Accepted date:

19 December 2013 14 May 2014 27 May 2014

Please cite this article as: A.S. Ramos, A.J. Cavaleiro, M.T. Vieira, J. Morgiel, G. Safran, Thermal stability of nanoscale metallic multilayers, Thin Solid Films (2014), doi: 10.1016/j.tsf.2014.05.065

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ACCEPTED MANUSCRIPT Thermal stability of nanoscale metallic multilayers

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A. S. Ramos1, A. J. Cavaleiro1, M. T. Vieira1, J. Morgiel2, G. Safran3

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CEMUC®, Departamento de Engenharia Mecânica, Universidade de Coimbra,

3030-788 Coimbra, Portugal 2

Reymonta 25, 30-059 Cracow, Poland 3

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Institute of Metallurgy and Materials Science, Polish Academy of Sciences,

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Research Institute for Technical Physics and Materials Science, Hungarian

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Academy of Sciences, H-1121 Budapest, Hungary

Abstract

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Metallic nanolayered thin films/foils, in particular Ni/Al multilayers, have been used to promote joining. The objective of this work is to evaluate the thermal stability of

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nanoscale metallic multilayers with potential for joining applications. Multilayers thin

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films with low (Ti/Al and Ni/Ti), medium (Ni/Al) and high (Pd/Al) enthalpies of exothermic reaction were prepared by dual cathode magnetron sputtering. Their thermal stability was studied by: i) differential scanning calorimetry combined with Xray diffraction (XRD), ii) in-situ XRD using cobalt radiation, and iii) in-situ transmission electron microscopy. It was possible to detect traces of intermetallic or amorphous phases in the as-deposited short period (bilayer thickness) multilayers,

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Corresponding author E-mail address: [email protected] Departamento de Engenharia Mecânica Pólo II, Universidade de Coimbra R. Luís Reis Santos 3030 Coimbra, Portugal Tel.: + 351 239790700; fax: + 351 239790701

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ACCEPTED MANUSCRIPT except for the Ti/Al films where no reaction products that might be formed during deposition were identified. For short periods (below 20 nm) the equilibrium phases

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are directly achieved upon annealing, whereas for higher periods intermediate trialuminide phases are present for Ti/Al and Ni/Al multilayers. The formation of B2-

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NiTi from Ni/Ti multilayers occurs without the formation of intermediate phases. On the contrary, for the Pd-Al system the formation of intermediate phases was never avoided.

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microparts/devices was demonstrated.

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The viability of nanoscale multilayers as “filler” materials for joining macro or

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Keywords: Multilayers; Ti/Al; Ni/Al; Ni/Ti; Pd/Al; Intermetallic; Sputtering; Thermal

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stability

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ACCEPTED MANUSCRIPT 1. Introduction Reactive multilayer (ML) thin films are composed of tens, hundreds, or thousands of

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alternating individual layers of reactants having a large negative enthalpy of mixing. The layered structure is repeated throughout the film with a modulation period ()

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consisting in the thickness of one double layer. The negative enthalpy of mixing of the constituents results in heat release once reaction is initiated. For certain systems and designs, reactive MLs exhibit fast exothermal reactions, attaining high

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temperatures. The reactions could be ignited by an external energy source such as

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an electric spark or a mechanical load [1]. After reaction is initiated, atomic mixing and diffusion take place generating heat at a very fast rate. Heat is conducted down

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the films, promoting atomic mixing and establishing a self-propagating chain reaction [1]. In nanoscale/nanocrystalline MLs heat can be generated faster, increasing the

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reaction velocity and enabling exothermic reactions to become self-sustained. The exothermic and self-propagating nature of reactive MLs make them promising for

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several applications, including joining [2]. In fact, reactive nanoscale ML foils and thin

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films have already found use as filler material, being particularly promising for dissimilar joining. Reactive MLs can be used as stand-alone heat sources or by using the energy released to melt surrounding solder or braze alloys. For some years now, the heat released by the exothermic reaction in nanolayered Ni/Al reactive ML foils is being used to melt braze alloys, promoting joining [3-5]. In addition, reactive MLs can be used to enhance the diffusion bonding process by taking advantage of their improved diffusivity/reactivity and exothermic character – reaction assisted diffusion bonding [6-8]. The feasibility of this procedure has already been demonstrated for similar and dissimilar joining using metallic ML thin films with nanometric modulation periods (bilayer thickness) [9-14]. The advantages of the reaction assisted diffusion 3

ACCEPTED MANUSCRIPT bonding process using reactive ML thin films as a “filler” material include: i) an “internal” energy source to aid on the diffusion bonding process, ii) lower diffusion

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bonding temperature/pressure or shorter holding times, iii) excellent filler/base material interface, iv) high quality and controllability of the filler material, and v) the

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possibility of providing an interface zone with an intermediate thermal expansion coefficient, avoiding cracks and fracture during dissimilar joining. Whatever the application envisaged, it is extremely important to study the thermal

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stability of the ML thin films and how they evolve with the temperature. The phase

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evolution of different metallic multilayer thin films has been studied by the authors using different techniques [8,15-18].

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This research concerns the phase evolution with temperature of Me1/Me2 (Me – metal) ML thin films with potential for joining applications. The thermal stability of ML

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thin films from different systems and with different modulation periods was studied by differential scanning calorimetry, hot X-ray diffraction and in-situ heating transmission

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electron microscopy. A thorough investigation regarding morphology, microstructure,

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phase composition and evolution with temperature should allow the most promising MLs for joining purposes to be identified.

2. Experimental Details 2.1 Multilayers’ Processing Me1/Me2 multilayer thin films were deposited by dual cathode magnetron sputtering at 0.3-0.4 Pa argon pressure using high purity metallic targets and a rotating substrates’ holder. The power density on each target was accordingly adjusted to achieve the desired equiatomic overall chemical composition (50 at.% Me1/50 at.% Me2). This atomic composition corresponds to different Me1/Me2 thickness ratios 4

ACCEPTED MANUSCRIPT depending on the Me1-Me2 system. Thin films with modulation periods ( - thickness of one layer of Me1 and one layer of Me2) between 4 and 40 nm were prepared by

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varying the rotation speed of the substrates’ holder, while 100 to 200 nm periods were achieved by alternately switching on and off the targets’ power supplies. The

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time that each target power supply was switched on was selected in order to have the selected individual layer thicknesses.

Prior to deposition, an ultimate vacuum pressure below 5x10-4 Pa was reached and

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the substrates’ surfaces were cleaned using an ion gun at 0.15 Pa argon pressure.

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To avoid heating during deposition, the substrates were on a thick copper block, which operated as a heat sink. Pd/Al, Ni/Al, Ti/Al and Ni/Ti multilayer thin films with

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nanometric modulation period and total thickness in the range 2-2.5 m were

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produced.

2.2 Multilayers’ Characterization

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The chemical composition of the multilayer thin films was evaluated by electron

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probe microanalysis (EPMA) using Cameca SX50 equipment operated at low voltages to ensure no influence of the substrate. Scanning electron microscopy (SEM) was used for the morphological characterization. The analyses were carried out at a 15 kV acceleration voltage in secondary and backscattered electron modes using a FEI QUANTA 400F high-resolution SEM equipped with a field emission gun. Microstructural analyses were conducted using a transmission electron microscope Tecnai G2 FEG 20 (200 kV) equipped with integrated Energy Dispersive Spectroscopy system. The nanostructure of the ML thin films was characterized using scanning-transmission High Angular Annular Dark Field. In-situ heating experiments were carried out using Philips CM20 transmission electron microscope. 5

ACCEPTED MANUSCRIPT In order to ensure adequate conditions of observation, namely reduction of image drift, in-situ transmission electron microscopy (TEM) experiments were performed in

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stages to allow drift stabilization. Thin foils for TEM were prepared using FEI Quanta 3D focused ion beam system equipped with Omniprobe lift-out system. Sample

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surface was coated with a layer of carbon allowing it to be separated from the platinum masking bar. Side trenches around the platinum bar were cut with 7 nA Ga+ current, while after removing thin foil from the surface and welding it to a copper grid

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a final polishing with 1, 0.5, 0.3 and 0.1 nA Ga+ currents was performed. The ML thin

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films were analysed by X-ray diffraction (XRD) using Co K radiation for phase identification. The XRD experiments were performed in a Philips diffractometer with

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Bragg-Brentano geometry and hot stage enclosed in a chamber under vacuum (optional), enabling XRD experiments at temperatures up to 1500ºC and controlled

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atmosphere. The high-temperature XRD thermal cycles were predefined taking into account the different Me1-Me2 systems. Differential scanning calorimetry (DSC)

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measurements were performed on free-standing films under purified argon

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atmosphere using Setaram G-DTA/DSC equipment. Heating rates () ranging from 10 to 45 ºC min-1 were used up to 800 ºC, while a 20 ºC min-1 cooling rate was used. The DSC experiments were accompanied with XRD analyses. Samples were heated just after the exotherms and then cooled to room-temperature (60 ºC min-1) and analysed by XRD.

3. Results and Discussion Metallic MLs from systems with different enthalpies of reaction were selected for this study (table 1). The reaction between the different metals is expected to result in the formation of equiatomic intermetallic compounds. High (Pd/Al), medium (Ni/Al) and 6

ACCEPTED MANUSCRIPT low (Ni/Ti and Ti/Al) energy ML thin films with nanometric modulation periods were prepared by magnetron sputtering. The EPMA results reveal near equiatomic overall

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chemical compositions, apart from some Ni/Al and Ni/Ti ML thin films that are slightly Al- and Ti-rich, respectively. Nevertheless, all the ML films under study fall in the

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Me1Me2 single phase domain of the respective equilibrium phase diagram. The Me1/Me2 multilayer thin films under study present a cross-section morphology type T according to Thornton’s model [21], which corresponds to a transition state

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between columnar and featureless morphologies. Cross-section SEM images of ML

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thin films where the alternating Me1 and Me2 layers can be distinguished ( ≥ 12 nm) are shown in figure 1 and 2. In figure 1 it is also possible to observe the globular

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surface morphology of a Pd/Al ML thin film with  = 30 nm. TEM analyses were conducted in order to observe the layered structure of short period MLs and check if

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reaction occurred during the deposition process. The bright-field (BF) TEM images (figure 3) reveal well-defined nanocrystalline layered structures. The geometry of the

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multilayer is to some extent compromised by waviness (figure 3) caused by the

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pseudo-columnar growth morphology, typical of films prepared by magnetron sputtering [22]. Wavy interface morphologies were observed and investigated in several vapour deposited Me1/Me2 MLs [23]. This waviness could influence the thermal stability of nanoscale MLs, in particular if full coherency is observed. The nanolayered structure is clear in scanning TEM mode (figure 4). Although the layered structure is observed, even for 4 nm period MLs (figure 4(b)), the possibility of reaction during the sputtering process must be considered. Me 1 and Me2 react at rather low temperatures (except Ti and Al) favouring the occurrence of intermixing and reaction during deposition. In fact, it was possible to confirm by XRD and TEM the formation of intermetallic and amorphous phases [7,8,15]. These phases form at 7

ACCEPTED MANUSCRIPT the interfaces being their contribution more evident for low periods, where the interfaces constitute a significant fraction of the films. In addition, in most as-

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deposited ML thin films it is possible to identify by XRD Me1 and Me2 diffraction peaks that became broader as the period decreases, indicating an increasingly

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nanocrystallinity and in some systems a predisposition for amorphization.

3.1 Phase Evolution

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The structural evolution of Ti/Al ML thin films with modulation periods from 4 up to

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200 nm was studied by DSC as complement of XRD [16]. The results confirmed that the period influences the phase evolution. Nevertheless, it should be noted that up to  = 200 nm it was always possible to form -TiAl as a major phase. Reducing the

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individual layer thickness (  20nm) leads directly to the equilibrium structure, while

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for higher periods an Al-rich intermediate phase was formed (TiAl3). The kinetic study revealed slightly lower activation energy values as the period decreases [16]. The

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phase evolution of Ti/Al ML films under slow heating was studied in detail by

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Rogachev and co-workers [2,24,25]. The DSC and time-resolved XRD results are consistent with those discussed above. In films with rather thick layers of reactants, the TiAl3 phase is initially formed followed by the formation of TiAl as the final product, whereas for  = 8 nm the reaction proceeds in one step [2,24,25]. The phase evolution with temperature of the more energetic Pd/Al and Ni/Al MLs was studied by hot-XRD using Co (K) radiation [8,18]. For each system, the thermal cycle, namely the maximum temperature, was defined in order to promote the formation of the monoaluminide intermetallic compound. As an example, a series of X-ray diffractograms of the 14 nm period Ni/Al MLs at different temperatures is shown in figure 5. The as-deposited film is nanocrystalline (Al + Ni), being difficult to 8

ACCEPTED MANUSCRIPT confirm by XRD the presence of the B2-NiAl intermetallic phase, since the Ni (111) and NiAl (110) diffraction peaks are almost coincident. As the temperature increases,

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the Al (111) peak tends to disappear while the peak at around 52º becomes more and more intense, which is attributed to the formation of the B2-NiAl phase. Similarly

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to the Ti/Al MLs, for 30 and 140 nm periods the formation of B2-NiAl during heating of Ni/Al ML thin films is preceded by the formation of Al-rich intermediate phases [18]. In Pd/Al ML thin films with nanometric modulation period the formation of

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intermediate phases was not avoided by reducing the individual layer thickness [8].

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Phase evolution in Ni/Al MLs has been extensively studied [e.g. 26-29]. Depending on the overall chemical composition, modulation period and processing history,

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different initial phases (NiAl3, Ni2Al9 and NiAl) have been reported and consequently different phase sequences were observed. In particular, the metastable Ni2Al9 phase

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not detected in the present study has been identified, either in Al-rich (60 at.% Al) Ni/Al MLs with periods down to 12.5 nm [28], either in MLs with equiatomic chemical

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composition and  between 30 and 85 nm [29]. On the contrary, Pd/Al films have

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received little attention and there are only a few papers regarding the phase formation and reaction during heating of Pd/Al ML films. Moreover, the films were never heated above 600ºC which makes not possible a comparison with the diffractograms at 700 and 800ºC, where the high temperature PdAl phase was detected [8]. Nevertheless, it is worth to mention the work of J. Braeuer et. al [30,31] who used high energetic Pd/Al MLs for reactive bonding at chip and wafer level. The phase evolution with temperature of Ni/Ti ML thin films was in-situ studied by hot-XRD using Co K radiation [15]. The formation of the B2-NiTi austenitic phase was detected at temperatures between 300 and 400ºC depending on the modulation period. The formation of B2-NiTi occurred in a single step; in this system the 9

ACCEPTED MANUSCRIPT formation of intermediate phases was never detected [15]. These results are corroborated by previous works with DSC combined with XRD [32,33]. By increasing

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the temperature above the equiatomic intermetallic formation, it was also possible to identify the presence of NiTi2 intermetallic phase. The formation of this brittle phase

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may be deleterious for joining purposes. The structural evolution with temperature of a Ni/Ti ML thin film with  = 25 nm was also analysed by in-situ TEM. The micrographs (figure 6) have shown that during heating the layered structure changes,

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as well as the corresponding selected area electron diffraction (SAED) pattern (20ºC

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versus 400ºC during 6 min). Extending the holding time at this temperature leads to grain growth (figure 6(e)). Analysis of the SAED patterns obtained at different stages

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of heating confirms the phase transition - Ni and Ti react just after 400ºC.The reaction seems to be concluded after a few minutes (the time required to switch the

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microscope BF observation mode to electron diffraction and vice versa, as well as the image acquisition time). As a result, Ni and Ti are no longer identifiable and a

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peak corresponding to (110) B2-NiTi appears. It is not possible to unequivocally

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confirm the presence of the austenitic B2-NiTi phase because the (100) superlattice peak is not observed, which may indicate a low ordering degree. Self-propagating reactions in Ni/Ti MLs were studied by Adams et al. [34]. Reacted MLs consist of one or more NixTiy phases depending on the modulation period; while the presence or absence of oxide phases is related with the reaction environment [34]. For small periods (< 29 nm) single phase austenite is formed with no evidence for NiTi 2 precipitates [34]. A summary of the phase evolution as a function of the ML system and modulation period is shown in table 2. According to the thermal stability analyses, the Ni/Al and Ti/Al ML thin films with intermediate modulation period seem to be the most 10

ACCEPTED MANUSCRIPT promising for using as reactive filler in diffusion bonding experiments. The formation of intermediate phases is avoided, and the effect of intermixing is less pronounced

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than for short periods (4 or 5 nm). For short periods the thickness of the intermixed layer is a non-negligible fraction of the ML period, reducing the amount of exothermic

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energy with a subsequent decrease of heat released and driving-force for selfpropagation. The heat released (higher for the Ni-Al system), the reaction temperature of the ML filler (higher for the Ti-Al system), and the base material(s)

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being joined will dictate the best selection for each case. Ni/Ti ML thin films constitute

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and filler materials are avoided.

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also a good option for joining NiTi shape memory alloys, disruptions between base

3.2 Reaction Assisted Diffusion Bonding

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Sound joints of TiAl to itself [10,11,13], TiAl to Inconel [12] and NiTi to Ti6Al4V [14] have been produced by reaction assisted diffusion bonding using Ti/Al, Ni/Al or Ni/Ti

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ML nanocrystalline thin films with a nanometric modulation period. TEM images of a

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NiTi/Ti6Al4V joint processed at 850ºC during 1h, under an applied load of 5 MPa, using a Ni/Ti ML thin film as filler material are presented in figure 7. The reaction assisted diffusion bonding of the ML coated NiTi and Ti6Al4V parts resulted in a sound joint without pores or cracks. Ni and Ti from the ML react, giving rise to intermetallic phases. Near equiaxed B2-NiTi grains can be identified close to the NiTi base material and at the centre of the joint interface, while columnar NiTi2 grains are observed on the Ti6Al4V side (figure 7(b)). The formation of this Ti-rich phase is favoured by the diffusion of Ti from the base material towards the filler material and of Ni in the opposite direction to beta-titanium phase.

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ACCEPTED MANUSCRIPT Similar and dissimilar joints have been systematically characterized by SEM and TEM confirming that it is possible to obtain sound joints (without pores or cracks) at

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less demanding bonding conditions even using low energy MLs such as Ti/Al and Ni/Ti [12-14]. Nanoindentation tests across the joints corroborate the SEM and TEM

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analyses, revealing higher hardness at the reaction zone, particularly at the interface with the base materials where Me1xMe2y (x  y) and Me1xMe2yMe3z intermetallic

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phases formed.

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Conclusions

The thermal stability of nanoscale metallic ML thin films with low (Ti/Al and Ni/Ti),

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medium (Ni/Al) and high (Pd/Al) heat exothermal reactions was evaluated. As the modulation period decreases, the occurrence of reaction cannot be avoided in the

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as-deposited Pd/Al, Ni/Al and Ni/Ti ML thin films. Heating the ML thin films leads to disruption of the layered structure by interdiffusion of Me1 and Me2, followed by

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chemical reaction to form intermetallic compounds. In agreement with the MLs

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overall chemical composition, it was always possible to obtain the desired equiatomic Me1Me2 intermetallic phase. Nevertheless, the phase evolution with temperature depends on the ML system and modulation period. Ti/Al and Ni/Al ML thin films with short periods (below 20 nm) directly give rise to equilibrium monoaluminide phases, whereas for higher periods intermediate trialuminide phases are formed. Whatever the modulation period the formation of B2-NiTi from Ni/Ti multilayers occurs without the formation of intermediate phases, while for the Pd-Al system the formation of intermediate phases was never avoided. Metallic ML thin films have been used to enhance the diffusion bonding process of advanced materials. In order to use these nanoscale MLs for joining applications, the 12

ACCEPTED MANUSCRIPT following guidelines should be considered: ML systems must have sufficient reaction energy to assist the joining process; the reaction should proceed without the

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formation of intermediate phases and therefore without energy losses; the modulation period has to be chosen in order to minimize the effect of intermixing;

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unsuitable interactions with the materials to be joined must be avoided.

Acknowledgements

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This research was supported by FEDER funds through the program COMPETE –

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Programa Operacional Factores de Competitividade – and by national funds through FCT – Fundação para a Ciência e a Tecnologia – under the project PEst-

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C/EME/UI0285/2013 and by FEDER/FCT funds through the project PTDC-EMETME-100990-2008 and through the Grant SFRH/BD/68354/2010. This research was

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also sponsored by FEDER funds through the program COMPETE – Programa Operacional Factores de Competitividade under project CENTRO-07-0224-FEDER-

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002001 “MT4MOBI - Materials and Technologies for Greener Manufacturing &

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Products Applied to Mobility”. The microstructure characterization was supported by Polish National Science Centre (NCN) through project DEC-2012/05/B/ST8/01794.

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[31] J. Braeuer, J. Besser, E. Tomoscheit, D. Klimm, S. Anbumani, M. Wiemer, T. Gessner, Investigation of different nano scale material systems for reactive

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wafer bonding, ECS Transactions 50 (2012) 241-251.

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[32] T. Lehnert, H. Grimmer, P. Boni, M. Horisberger, R. Gotthardt, Characterization of shape-memory alloy thin films made up from sputter-deposited Ni/Ti

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multilayers, Acta Mater. 48 (2000) 4065-4071. [33] R. Gupta, M. Gupta, S. Kulkarni, S. Kharrazi, A. Gupta, S. Chaudhari, Thermal

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stability of nanometer range Ti/Ni multilayers, Thin Solid Films 515 (2006) 22132219.

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[34] D.P. Adams, M.A. Rodriguez, J.P. McDonald, M.M. Bai, E. Jones Jr, L. Brewer,

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J.J. Moore, Reactive Ni/Ti nanolaminates, J. Appl. Phys. 106 (2009) 093505.

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ACCEPTED MANUSCRIPT Figure Captions

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Figure 1 – Cross-section SEM micrograph of a Pd/Al ML thin film with  = 30 nm.

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Grey and dark layers correspond to (Pd) and (Al), respectively.

Figure 2 – Cross-section SEM micrograph of a Ni/Ti ML thin film with  = 70 nm.

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Grey and dark layers correspond to (Ni) and (Ti), respectively.

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Figure 3 – Cross-section BF TEM micrographs of Me1/Al multilayer thin films with (a) Me1=Ti,  = 20 nm and (b) Me1= Ni,  = 30 nm. Grey and dark layers correspond to

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(Al) and (Me1), respectively.

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Figure 4 – Cross-section STEM micrographs of Ni/Al ML thin films with (a)  = 30 nm

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and (b)  = 5 nm. Grey and dark layers correspond to (Ni) and (Al), respectively.

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Figure 5 – Hot-stage XRD diffractograms of a Ni/Al ML thin film with  = 14 nm. The stainless steel substrate most intense diffractions are labeled with “S”.

Figure 6 – (a), (c), (e) BF TEM micrographs and (b), (d), (f) corresponding SAED patterns of a Ni/Ti multilayer thin film with  = 25 nm at (a), (b) room-temperature (c), (d) 400ºC/6min and (e), (f) 400ºC/20min. The diffraction rings corresponding to the silicon substrate are labeled with “S”.

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ACCEPTED MANUSCRIPT Figure 7 – BF TEM micrographs of a NiTi/Ti6Al4V joint processed at 850ºC / 5 MPa / 1h using a Ni/Ti multilayer with  = 25 nm. (a) Joint interface. (b) Magnification of the

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joint interface on the Ti6Al4V side.

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ACCEPTED MANUSCRIPT Table Captions

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Table 1 – Enthalpy of reaction (H0) and adiabatic reaction temperature (Tad) of

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bimetallic systems [1,19,20].

Table 2 – Phase evolution with temperature of Me1/Me2 multilayer thin films

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[8,15,16,18].

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Figure 1

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Figure 2

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Figure 3a

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Figure 3b

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Figure 4a

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Figure 4b

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Figure 5

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Figure 6a

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Figure 6c

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Figure 6e

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Figure 6f

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Figure 7a

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Figure 7b

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ACCEPTED MANUSCRIPT Table 1 – Enthalpy of reaction (H0) and adiabatic reaction temperature (Tad) of

Ti/Al

-92

-59

-36

2380

1639

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Tad (ºC)

Ni/Al

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H0 (kJ mol-1)

Pd/Al

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bimetallic systems [1,19,20].

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Ni/Ti -33 1355

ACCEPTED MANUSCRIPT Table 2 – Phase evolution with temperature of Me1/Me2 multilayer thin films

Tf2 (ºC)

Period (nm)

Phase Transformation Sequence

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Me1/Me 2 System

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[8,15,16,18].

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575-650 Al + Ti → Al + Ti + γ -TiAl → γ-TiAl + 2-Ti3Al (traces)

Ti/Al

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625-680 Al + Ti → Ti + TiAl3 → γ -TiAl+ 2-Ti3Al (traces)

Ni/Al

5, 14

150-225 Al + Ni → Al + Ni + B2-NiAl → B2-NiAl

Ni/Al

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175-225 Al + Ni → Ni + NiAl3 → Ni + NiAl3 +B2-NiAl → B2-NiAl

Ni/Al

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325-500

Al + Ni → Al + Ni + NiAl3 → NiAl3 + Ni2Al3 → Ni + Ni2Al3 + B2-NiAl → B2-NiAl

Al/Pd

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700-800

Al + Pd + m-AlPd → Pd + m-AlPd + Al3Pd2 → Al3Pd2 + LT-AlPd → LT-AlPd + β-AlPd → β-AlPd

Al/Pd

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700-800

Al/Pd

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Ni/Ti

5, 12, 25, 70

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Ti/Al

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Al + Pd + m-AlPd → Pd + m-AlPd + Al3Pd2 → Al3Pd2 + LT-AlPd → LT-AlPd + β-AlPd → β-AlPd Al + Pd → Pd + (m-AlPd) + Al3Pd2 → Pd + Al3Pd2+ 700-800 LT-AlPd → Al3Pd2 + LT-AlPd → LT-AlPd + β-AlPd → β-AlPd

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300-400

Ni + Ti → B2-NiTi → B2-NiTi + NiTi2

m-AlPd – Metastable phase

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LT AlPd – Low temperature phase

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Tf – Temperature range corresponding to the formation of the Me1Me2 equiatomic intermetallic

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ACCEPTED MANUSCRIPT Highlights 

Me1 and Me2 (Me – metal) alternated nanolayers deposited by magnetron

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sputtering. Reactive Me1/Me2 multilayer thin films with nanometric modulation period.

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By heat treatment the films always evolve to the equilibrium intermetallic

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phase. 

For some Me1-Me2 systems and periods, the formation of intermediate phases occurs.

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Me1/Me2 multilayer thin films can be used as filler materials to enhance

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joining.

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

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