58 (2007) 555–562

Materials Characterization

Characterization of interface of Al2O3–304 stainless steel braze joint

Abhijit Kar, Sudipta Mandal, K. Venkateswarlu, Ajoy Kumar Ray ⁎

Material Science and Technology Division, National Metallurgical Laboratory (CSIR), Jamshedpur-831 007, India

Received 4 September 2006; received in revised form 4 December 2006; accepted 4 December 2006

Abstract

Sintered Al2O3 is brazed with 304 stainless steel (SS) using 97(Ag28Cu)3Ti active filler alloy at 900 and 1000 °C. Theinterfaces developed during brazing have been systematically characterized and the results show that the thickness of the interfacesis 15–20 μm and 10–15 μm on SS side and Al2O3 side respectively. The presence of TiO, Cu3Ti3O and FeTi phases at the Al2O3

interface and FeTi, and Fe35Cr13Ni3Ti7 phases at the SS interface have been confirmed through XRD and TEM studies.Microhardness analysis across the brazed interfaces suggest that the SS interface (300–600 Hv) is harder than that of the SSsubstrate (200–250 Hv) whereas alumina interface (550–900 Hv) is softer than that of the alumina substrate (∼1900 Hv). It isconcluded that the various new phases that are formed during brazing are responsible for the variation in the hardness values.© 2006 Elsevier Inc. All rights reserved.

Keywords: Ceramic brazing; Ag–Cu–Ti filler alloy; Interface; Phases; Characterization

1. Introduction

In view of the superior physical, chemical and me-chanical properties of the ceramics, they are widely usedin the electronics, aerospace, nuclear, automobile, cut-ting tool and biomaterials sectors [1,2]. To utilize thehigh temperature properties of ceramics, it is oftenrequired to join the ceramics with metallic parts. Asignificant difference in thermal and mechanical prop-erties of ceramics and metals makes it difficult to obtainceramics–metal joints with adequate mechanical integ-rity [3]. The joining of ceramics for high temperatureapplications is difficult due to the poor wettability of theceramic substrates by the conventional brazing alloy[4,5]. Therefore, it is necessary to develop a reaction

⁎ Corresponding author. Tel.: +91 657 227 1709; fax: +91 657 2270527.

E-mail address: raykay50@yahoo.co.in (A.K. Ray).

1044-5803/$ - see front matter © 2006 Elsevier Inc. All rights reserved.doi:10.1016/j.matchar.2006.12.001

layer at the ceramic surfaces to increase their wettingbehaviour with the liquid filler alloy [6,7]. One suchmethod is the moly-manganese metallization processdeveloped for the joining of ceramics [8] but the processhas restricted usage because of several stringent processparameters.

Active metal brazing has been considered as anotherway of joining, which is extensively used, in the recentpast. Addition of chemically active elements such as Tito the filler alloy improves its wettability to ceramicsubstrate [9]. It has been observed that the filler alloycontaining an active metal like Ti, can join ceramicswithout metallization, because the active metal react tochange the chemistry of the ceramic surfaces [10]. Thisprocess is cost-effective and carried out in a single step.In addition, the joining temperatures can be variedbetween 650 and 1300 °C as required.

The ability to produce a reliable ceramic–metal brazejoint is a key enabling technology for many advanced

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material applications. During joining of two heteroge-neous materials, many changes in microstructure andmechanical properties in the vicinity of joining regionare observed [11]. For any brazed component, theadmissible working temperature is dependent more onthe nature of the interface rather on the substrates. Theinterfacial microchemistry plays important roles indetermining a successful braze joint, thus it is necessaryto identify/know the reaction products formed at theinterface [12].

Since alumina is most effectively brazed at 900 and1000 °C through active brazing process [13], the brazingtemperatures for the present investigation have beenselected accordingly. The reaction products formed atthe interfaces during the brazing of Al2O3–304 SS havebeen characterized by SEM–EDX, EPMA, XRD andTEM–EDX analyses. An attempt has been made to

Fig. 1. SEM–EDX analyses of the active filler alloy (a) o

evaluate the relative microhardness of the interfaces,substrates and the filler alloy.

2. Experimental procedures

The starting materials used for the preparation ofactive filler alloy were Ag, Cu, and Ti with 99.9%, 99.9%and 99.5% purity respectively. The alloy, 97(Ag28Cu)3Ti was prepared by melting and casting route. Details ofthe alloy preparation were reported elsewhere [13,14].Alumina substrate was made by sintering A16SG (Alcoa,USA) powder at 1600 °C using electric resistance fur-nace. The holding time for sintering was 30 min at thehighest temperature. The density of the sintered aluminawas 3.78 g/cc. The sintered alumina was brazed to 304stainless steel (Fe18Cr8Ni2Mn). The substrates and thefiller alloys were cut by precision cut-off machine

ver all analysis, (b) bright phase and (c) dark phase.

Fig. 2. XRD analyses around the Al2O3 interface of Al2O3–304 SS braze joint; (a–c) represent mostly Al2O3 substrate, (d and e) represent mostlyinterface along with Al2O3, Ag and Cu, (f) represents mostly the residual filler alloy.

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(MECATOME P100, Presi, France) using a diamondwafer blade. The dimensions of these materials used forjoining are as follows:

Al2O3 10(l)×8(b)×5(t) mm304 SS 10(l)×8(b)×5(t) mmFiller alloy 9(l)×7(b)×0.4(t) mm

Fig. 3. X-ray diffraction pattern of Al2O3–3

Roughness (Ra) of the Al2O3 is 0.45 μmRoughness (Ra) of the 304 SS is 0.16 μm.

Cut substrates and the alloys were polished andcleaned ultrasonically using ethanol and acetone. Theroughness of the cleaned substrates was measured byprofilometer (Taylor Hobson precision, Taylsurf Series

04 SS brazed joint at the SS interface.

Fig. 4. (a) Backscattered electron image (BSEI) of Al2O3–304 SSbraze joint and (b) EPMA dot map for Ti.

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2). The filler alloy was sandwiched between the aluminaand the stainless steel substrates. The assembly wasplaced in a graphite resistance furnace (ASTRO,Thermal Inc., USA). A fixed load of ∼1 kg was keptover the sample to keep the assembly aligned duringbrazing. This enables a proper contact between thesurface of the substrates and the filler alloy [5,15–17].This load also facilitates the interfacial reaction [18].The furnace was heated at the rate of 6 °C/min. till itreached 900 °C and kept for 15 min at that temperature.The cooling rate was maintained at 3 °C/min till itreached to 500 °C and thereafter furnace cooled. Thesame brazing cycle was maintained for another separateexperiment but the maximum brazing temperaturemaintained was at 1000 °C. Before heating, the furnacewas purged with argon gas (IOLAR I), and through outthe brazing cycle 30–50 mTorr Ar pressure wasmaintained by rotary pump.

The brazed samples were cut in order to get a freshsurface for characterization purpose. The cut sampleswere metallographically polished with 0.5 μm diamondpaste and subjected to carbon coating for SEM (JEOL-840A, Japan) and EPMA (JXA — 8600 M, Jeol, Japan)analyses. Line profile quantitative elemental analyseswere carried out by EPMA across the brazed interface.Phases present at the interfaces were identified by XRD(Seifert, PTS 3003, Germany). The Al2O3 was removedstep by step to reveal the interface adjacent to aluminasubstrate. Similarly, the stainless steel was also removedfor analyzing the interface adjacent to stainless steel. Afterremoval of each layer XRD analysis was carried out. Thesample prepared for TEM analysis [16] was ultrasonicallycleaned with acetone and was analyzed by TransmissionElectronMicroscope (CM200 Philips, Netherlands, fittedwith EDAX — DX 4) for microstructures and identifi-cation of the phases present at the interface.

Microhardness measurements were carried out start-ing from stainless steel to alumina covering both theinterfaces and intermediate residual filler alloy. Themicrohardness tests were performed using Vickers hard-ness tester (Leitz, RZD-DO, Germany). A load of 50 gand 100 g was applied on SS/filler alloy and ceramicsurfaces respectively. An average hardness value hasbeen reported by taking five readings in each case.

3. Results and discussion

3.1. SEM–EDX analysis of 97(Ag28Cu)3Ti active filleralloy

The SEM–EDX analysis of the filler alloy revealedthat most of the titanium was associated with copper rich

dark phase (Fig. 1). A low quantity of titanium wasfound in the silver rich bright phase. This is due to theformation of some intermetallic phases, e.g. Cu3Ti[17,18]. Titanium has a solubility of 67 wt.% at 1150 °Cin copper whereas it is only 3% in silver. In the presenceof silver, the tendency of titanium to associate with themolten alloy gets reduced and thereby increasing theactivity of Ti in the Ag–Cu–Ti alloy [19,20]. Smallamount of (∼3 wt.%) titanium is more than adequatelyactive to ensure wetting of even highly refractoryceramics such as alumina, silicon carbide, silicon nitrideor graphite [21].

3.2. XRD analysis

Fig. 2 shows the XRD patterns of braze joint at thealumina interface. Fig. 2 (a–c) represents mostly Al2O3

substrate. Fig. 2 (d, e) represents reaction products alongwith Al2O3, Ag, Cu. Fig. 2 (f) represents residual filleralloy (mostly Ag–Cu after the diffusion of Ti). The

Fig. 5. EPMA line profile analyses of the Al2O3–304 SS brazed interface (a) SS interface and (b) Al2O3 interface.

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interface adjacent to alumina was very thin and theconcentration of reaction products was also less inquantity, as a result, intensity of the XRD peaksobserved is small. The small peaks that are observedin Fig. 2 (d, e) are identified as TiO, Cu3Ti3O at thealumina interface. Further confirmation of these phasesare carried out by TEM analysis.

The interface adjacent to SS layer was large and XRDpeaks of the reaction products are prominent enough toconfirm the different reaction products at the interface.

Fig. 3 shows SS interface consists of FeTi and Fe35Cr13-Ni3Ti7 phases. Since the XRD peaks of both FeTi andNiTi are very close to each other, it is difficult todifferentiate them. However, the greater stability of FeTiphase as compared to the metastable NiTi phase andthe higher concentration of Fe (∼67 wt.%) than Ni(∼10 wt.%) in the SS substrate, FeTi has been consideredas the most probable phase. Ti from the Ag–Cu–Ti filleralloy has been diffused towards SS side and reacted withthe constituents present in the SS to form phases like FeTi

Fig. 6. TEM diffraction pattern, corresponding micrograph and EDX analysis of alumina interface; representing (a) TiO phase and (b) Cu3Ti3O phases.

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and Fe35Cr13Ni3Ti7. Formation of these phases in generalis responsible for bonding between SS and filler alloy.

3.3. Electron probemicroanalysis (EPMA) of the braze joint

The backscattered electron images (BSEI) and theX-ray dot maps of transverse section (Fig. 4) of theAl2O3–SS brazed with 97(Ag28Cu)3Ti active filleralloys show that there are distinct and continuousreaction layers between the filler alloy and Al2O3 aswell as between the filler alloy and SS. It is observedthat major part of the Ti from the filler alloy has beendiffused towards the SS substrate. As a result, the widthof the SS interface (15–20 μm) is bigger than that ofalumina interface (10–15 μm) (Fig. 4b).

Electron probe microanalyses (EPMA) of theAl2O3–SS brazed at 900 °C are shown in Fig. 5 (a, b).It has been observed that the joining process is due to theinter-diffusion of Ti, Ag and Cu towards the SS substrateand Fe, Ni, Cr towards the filler alloy leads to theformation of the SS interface, Fig. 5 (a). Similarly, thediffusion of Ti, Ag and Cu from the filler alloy towardsAl2O3 and Al from Al2O3 towards the filler alloy formsthe Al2O3 interface, Fig. 5 (b). It has been reportedelsewhere that diffusion of Ti is the key factor for theformation of the interfaces [22]. Therefore, the thicknessof the interfaces has been measured on the basis of thediffused titanium at the interfaces and found to be more orless the same as in the case of EPMA dot map analysis.

Fig. 7. Line profile microhardness results, measured across the inte

Yang et al. [23] observed that the thickness of the reactionlayer is a function of processing time and temperature fora particular brazing system.

3.4. TEM analysis

Since the XRD peak intensities for the reactionproducts are very small, the phases are further confirmedby TEM analysis. Fig. 6 (a) exhibits the SAD pattern forthe TiO and Fig. 6 (b) represents the SAD pattern for theCu3Ti3O phases [24]. The zone axis of TiO and Cu3Ti3Ois found to be [011] and [001] respectively. The cor-responding detailed indexing of the planes is depicted inthe insets of the respective figures (Fig. 6a and b).

3.5. Microhardness analysis

In general, microhardness is a function of orientationof basal plane and the indentation load. This orientationand load dependency can be clearly observed when theexperiment is carried out at low load, although at higherload microhardness becomes load and orientationindependent. This is also true for hard materials likealumina. Kaji et al. observed that in case of brittlematerials, using higher load causes microcracks sur-rounding the indent thus produces unsymmetricalindentation; hence results are spurious [25,26]. In thepresent study it has been observed that the use of 50 gand 100 g load on SS-filler and filler-alumina interface

rfaces, starting from the Al2O3 substrate to 304 SS substrate.

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respectively exhibits the best results. Line profile micro-hardness value of various zones of brazed joint is shownin Fig. 7. The average hardness value of the SS substrateis 230 Hv. While taking the hardness values from SSto interface, a Gaussian nature of the hardness profile(Fig. 7) across the SS interface (∼15 μm) is observed.The hardness values in this region are in the range of300–600 Hv. From XRD study, intermetallic phasesformed during brazing have been identified as FeTi andFe35Cr13Ni3Ti7. The microhardness results obtained atthe SS substrate and SS interface also suggest that theinterface is harder than the SS and this is due to thepresence of such intermetallics. On the other hand, thehardness values of alumina interface are in the range of550–900 Hv and that of Al2O3 substrate is ∼1900 Hv.

Hardness of the residual filler alloy, exists betweenthe two interfaces has been found to be ∼125 Hv(Fig. 7). It is also observed that the residual filler alloyjust adjacent to the SS interface has the hardness valuesof ∼140 Hv that is little higher than 125 Hv. Thereaction products are responsible for the hardness ofboth the interfaces.

4. Conclusions

1. A 304 SS has been successfully brazed to sinteredalumina at two different temperatures i.e. 900 and1000 °C, using 97(Ag28Cu)3Ti active filler alloy.

2. Lesser amount of reaction products formed at thealumina interface caused a thinner alumina interface(10–15 μm) than SS interface (15–20 μm).

3. Phases like TiO and Cu3Ti3O at alumina interface,whereas FeTi and Fe35Cr13Ni3Ti7 at SS interfacehave been confirmed.

4. Presence of various intermetallic phases at the inter-face is responsible for change in the microhardnessresults.

Acknowledgements

The authors would like to thank Prof. S.P. Mehrotra,Director, NML for his kind permission to publish thiswork. The authors are also thankful to the Department of

Science and Technology and Aeronautics Research andDevelopment Board, Government of India for theirpartial financial support to carry out this work.

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