Materials Science and Engineering A 424 (2006) 163–173

An investigation of fusion zone microstructures in electronbeam welding of copper–stainless steel

I. Magnabosco a, P. Ferro a,∗, F. Bonollo a, L. Arnberg b

a Department of Management and Engineering, University of Padova, Stradella S. Nicola, 3 I-36100 Vicenza, Italyb Department of Materials Science and Engineering, Norwegian University of Science and Technology, 7491, Trondheim, Norway

Received 1 October 2005; received in revised form 1 February 2006; accepted 1 March 2006

Abstract

The article presents a study of three different welded joints produced by electron beam welding dissimilar materials. The junctions were obtainedbetween copper plates and three different austenitic stainless steel plates. Different welding parameters were used according to the differentthicknesses of the samples. Morphological, microstructural and mechanical (micro-hardness test) analyses of the weld bead were carried out.

The results showed complex heterogeneous fusion zone microstructures characterized both by rapid cooling and poor mixing of the materialswhich contain main elements which are mutually insoluble. Some defects such as porosity and microfissures were also found. They are mainlyd©

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eywords: Dissimilar metal welding; Electron beam welding; Fusion zone microstructure; Stainless steels; Copper

. Introduction

Welding of dissimilar metals has been the object of inves-igations for many years; their growing importance is justifiedy their technical and economic potentials. Dissimilar metalsr alloys are welded in order to gain flexibility but this practiceften yields problems that negatively affect the performance ofhe weld.

A question of primary importance is the evaluation of thentrinsic properties of the materials to use in a dissimilar weld-ng. Many researchers [1–7] studied the effects of fusion weldingf materials characterized by different electrochemical, thermal,ptical and mechanical properties. Dissimilar welds carried outith alloys having widely different thermal expansion coeffi-

ients, fail either during fusion zone (FZ) solidification or duringhe service life, due to thermal fatigue [1]. The welding of two

etals (or alloys) having different melting temperatures or ther-al conductivities is complicated by the fact that one materialill melt prior to the other one [1]. Particular welding parame-

ers, such as the heat source position, have to be optimised inrder to produce a symmetric bead shape. The metallurgical

compatibility must also be taken into account in order to preventthe formation of brittle phases and the segregations of high or lowmelting point phases inside the bead. A general rule which canbe used in planning for dissimilar welding is that the mechanicaland corrosion resistance of the joint has to be at least compa-rable to that of the metal characterized by the lower properties;furthermore, it is important to reduce the presence of intermetal-lic phases in order to guarantee good mechanical and chemicalproperties.

The choice of the welding process is essential in order to con-trol the above-mentioned phenomena. The use of high-powerdensity heating sources, such as electron beam or laser beam,minimizes some of these problems thanks to the greater controlof shape and dimensions of the molten region and beam align-ment; the high power density permits one to reach immediatelythe fusion of both metals even if they have very different meltingtemperatures.

Recently, Mai and Spowage [4] studied a laser weldingprocess between copper and tool steel using a 350 W pulsedNd:YAG laser. They found that a major problem was hot crack-ing in the heat-affected zone (HAZ) of the steel plate induced

∗ Corresponding author. Tel.: +39 0444 998743; fax: +39 0444 998888.E-mail address: ferro@gest.unipd.it (P. Ferro).

by the diffusion of copper into the grain boundaries of the un-melted steel. Butt joints were thus created by focusing the laserbeam 0.2 mm into the steel. The results showed a very limi-ted (<2%) amount of copper dissolved in the molten steel, the

921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.03.096

164 I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173

absence of micro-cracks in the FZ–HAZ interface and a dendriticmicrostructure with the presence of martensite in the FZ.

Studying a copper to steel electron beam weld, Wei et al.[6] observed a homogeneous FZ with a chemical compositioncharacterized by gradual variations in the concentrations of ironand copper. Even though the melting temperatures of copper andiron differ from one another by about 450 ◦C, the differences indepths of the FZ or in the solid–liquid interfaces were negligible.This was attributed to the high temperature gradient induced bythe electron beam welding process.

Finally, Tosto et al. [7] carried out dissimilar butt weldedjoints between copper and AISI 304L stainless steel. The sam-ples were 25 mm in thickness and welded using the electronbeam welding (EBW) process. The authors did not find defectssuch as porosity or microfissuring in FZ and HAZ even if theyemphasize that some problems arose during service life due tothe supposed instability of the observed microstructures.

It should be pointed out that in the past only a few researchershave investigated dissimilar copper–steel welds. Specific infor-mation on process and geometrical parameters and their effectson microstructure and defects of the joint have, therefore, notyet been fully established.

The present work is focussed on the morphological,microstructural and mechanical analyses of some dissimilarelectron beam welded joints in copper and austenitic stainlesssteel in order to evaluate their potential in view of industrialani

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Table 2Operating and geometrical welding parameters

Parameters Joint 1 Joint 2 Joint 3

Cathode current (A) 2.35 2.35 2.35Voltage (kV) 150 150 150Beam current (mA) 90 120 130Gun vacuum (mbar) 10−5 10−5 10−5

Chamber vacuum (mbar) 10−4 10−4 10−4

Beam focus point (mm) 0 0 0Welding speed (mm/s) 6 6 6Gun-specimen distance (mm) 320 450 575Welding pass (#) 1 2 2Plates thickness (mm) 29 (steel)–33 (copper) 50 71

according to the thickness of the samples. Table 2 summarizesthe different parameters used in the present investigation. Theelectron beam was centred on the contact surfaces of the plates.The process parameters were chosen on the basis of industrialexperience. Due to the high thickness of the specimens 2 and 3,a two pass weld was chosen in order to limit the heat input.

The microstructure of the FZ and HAZ was studied by opti-cal and scanning electron microscopy. The chemical etchant forcopper was a diluted solution of ammonium persulphate. Theelectrolytic etch for steel was a solution of 10 g oxalic acid in100 mL H2O with a voltage 6 V and a cathode current of 40 mA,applied for 5 min.

Image analysis was used to obtain a quantitative evaluationof the area fraction of both ferrous and copper matrices in theFZ. This analysis was done at different positions in the seam.

A detailed investigation of the distribution of chemical ele-ments in the fusion zone was carried out by means of energydispersive X-ray spectroscopy, EDS, in an environmental scan-ning electron microscope. Point counting and line scan intensityprofiles for Fe(K), Cr(K), Ni(K), Cu(K) were obtained in orderto determine the elemental distributions across the fusion zone.

Vickers microhardness profiles were performed across thesection of the weld bead in order to identify the effects ofmicrostructural heterogeneities.

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pplications in the field of heat transfer devices. Different thick-ess, welding parameters and chemical compositions were takennto account.

. Materials and methods

Three electron beam weldments, carried out by the “E. Zanon.p.A.” industries in Schio (Italy), were studied. The first jointas obtained between a plate of copper ETP (UNS C11000

8,9]) and a plate of AISI 304L stainless steel. The second jointas produced by joining copper ETP (UNS C11000 [8,9]) toISI 304 stainless steel. The third joint was obtained by join-

ng a copper alloy (UNS C18100 [8,9]) to AISI 316L stainlessteel. Table 1 shows the measured compositions of the steellates and the typical compositions of copper alloys used in thisork.The alloys were electron-beam welded using different

perating parameters, optimised in order to obtain a full pene-ration. In particular, single- or double-pass welding was used,

able 1hemical compositions of the materials used in the tests (wt.%)

C Mn Si S P

04L 0.012 1.580 0.480 0.003 0.0204 0.065 1.670 0.632 0.015 0.0216L 0.030 1.060 0.620 0.015 0.02NS C11000 [9] Electrolytic tough pitch copper (ISO Cu-ETP) has a mi

99.90% with 200–600 ppm (0.02–0.06%) oxygen preseCu

NS C18100 [9] Balance

. Results

.1. Joint 1: Cu–AISI 304L

The first sample was obtained by a single-pass welding of twolates, one of copper and the other of AISI 304L; the geometrical

Cr Ni Mo Cu Fe

18.473 11.205 0.412 0.365 Balance17.791 8.301 0.564 0.532 Balance16.910 11.334 2.080 0.150 Balance

m copper ofcuprous oxide

Zr Mg

1.2 0.05–0.3 0.03–0.06

I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173 165

Fig. 1. Weld bead of the Cu–AISI 304L joint 1. Left: FZ–copper and right: FZ–AISI 304L interfaces.

Fig. 2. Micrographs of the FZ of joint 1: Cu–AISI 304L. (a) Steel globules embedded in a copper matrix; (b) steel matrix containing a copper globule.

characteristics of the bead are described in Table 3. Fig. 1 shows aschematic profile of the FZ and two micrographs of the FZ–HAZinterfaces.

Fig. 2a shows the microstructure in the top region of the beadcharacterized by steel globules and dendrites embedded in thecopper matrix. On the other hand, Fig. 2b shows the microstruc-ture in the centre of the FZ; a steel matrix containing a copperglobule in which fine globules of steel can be observed. It wasfound that both the dendrite size and the amount of dendritesdecreased as one moves from the top of the joint. The secondarydendrite arm spacing had a mean value of 1.6 m in the top zoneof the FZ and 1.4 �m in the centre.

Metallographic inspections of joint 1 did not reveal defectssuch as microcracks or porosity in the as-welded specimen;however, some copper penetrations were detected in austeniticgrain boundaries at the interface between steel and the FZ(Fig. 3).

Due to the colour contrast obtained in the optical micrographsof the FZ, it was easy to identify the two different phases, oneof them rich in Fe–Ni–Cr and the other in Cu. Image analysiswas done at three positions in the FZ, at a magnification of 50×.The obtained results are summarized in Table 4 from which onecan calculate a average area fraction of copper rich phase in theFZ of about 57%.

TG

MMD

Fig. 4a shows dark globules embedded in a light matrix whileFig. 4b displays a globule detail. Fig. 5 shows the correspondingEDS spectra. Point A of Fig. 4 is related to light matrix and,point B of Fig. 4 to dark globules whereas point C shows darkfeatures in the light matrix and point D lighter features in thedark globule.

Zone A is rich in copper with low concentrations of iron andchromium (Fig. 5A). The darker zone C, inside the matrix A,shows a higher concentration of copper and a lower presence ofthe other elements (Fig. 5C). Zone B is mainly characterized byiron and chromium (Fig. 5B) with a low concentration of copperwhich increases in the lighter D areas (Fig. 5D).

Fig. 6 shows elemental line scans across the fusion zonesand heat affected zones. It can be seen that the iron- chromium-

F(

able 3eometrical parameters of the FZ of joint 1: Cu–AISI 304L welding

aximum depth (mm) 30.13aximum width (mm) 4.16epth to width ratio 7.24

ig. 3. Micrograph showing copper penetrations at the interface between steel

base material) and FZ of joint 1: Cu–AISI 304L.

166 I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173

Table 4Quantitative evaluation of the percentage of both ferrous matrix and copper matrix in the FZ of joint 1: Cu–AISI 304L welding

Zone Copper side Centre Steel side Total

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

Zone 1 59 41 80 20 63 37 67.3 32.7Zone 2 58 42 48 52 41 59 49 51Zone 3 60 40 55 45 46 54 53.6 46.4

Fig. 4. SEM micrographs of the FZ of joint 1: Cu–AISI 304L. (a) Dark globules in the light matrix; (b) detail of a globule.

and nickel-concentration profiles coincide and that the elemen-tal distributions vary significantly over the zone showing avery heterogeneous composition. The microhardness distribu-tion measured on three different depths of the FZ is shown inFig. 7; even if such measures confirm the heterogeneity of themicrostructure, one can conclude that the mechanical strengthof the bead is equal or greater than that of the weaker basematerial.

3.2. Joint 2: Cu–AISI 304

The second sample analysed was a double-pass welded jointobtained by welding together a plate of copper (type ETP)and one of type AISI 304 stainless steel. The geometricalcharacteristics of the bead are described in Table 5, where num-bers 1 and 2 correspond to the first and second welding passcarried out on opposite sides of the sample.

ts A, B

Fig. 5. EDS spectra of poin , C and D zones of Fig. 4.

I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173 167

Fig. 6. Line profiles across the fusion zone of joint 1: Cu–AISI 304L.

Fig. 7. Microhardness profiles measured on three different depths of the beadof joint 1: Cu–AISI 304L.

Table 5Geometrical parameters of the FZ of joint 2: Cu–AISI 304 welding

Maximum depth 1 (mm) 32.5Maximum depth 2 (mm) 23.4Maximum width 1 (mm) 7.1Maximum width 2 (mm) 6.8Depth to width ratio 1 4.6Depth to width ratio 2 3.4

Fig. 8 shows a schematic profile (centre) and two micrographsof FZ–base metal interfaces.

The SEM micrographs of Fig. 9 show an intergranularmicrocrack at the interface between the FZ and the steel plate(Fig. 9a) and a shrinkage pore inside a copper globule (Fig. 9b).Fig. 10a shows the microstructure of the bottom region of thebead characterized by steel globules embedded in a copper-richmatrix, in which dendritic structures can be seen. Fig. 10bshows the steel matrix containing a globule of copper, in whichfine globules of steel have solidified. This last micrographis obtained from the centre zone of the bead where the twowelding passes overlap. The secondary dendrite arm spacinghas a mean value of 1.5 and 1.7 �m in the bottom and top zoneof the bead, respectively.

The area fraction of the ferrous- and copper-rich matrices aresummarized in Table 6. The image analysis was carried out infive zones of the seam at a magnification of 100×. The averagearea fraction of the copper-rich phase in the FZ is 62%.

Microchemical analysis of this weld showed results similarto that of joint 1.

Microhardness profiles measured on three different depths ofthe FZ are shown in Fig. 11. The microhardness mean value of

Fig. 8. Weld bead of joint 2: Cu–AISI 304. Left: F

Z–copper and right: FZ–AISI 304 interfaces.

168 I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173

Fig. 9. SEM micrographs of the FZ of joint 2: (a) microcrack at the interface between the FZ and steel; (b) porosity in a copper globule.

Fig. 10. Optical micrographs of the FZ of joint 2: Cu–AISI 304. (a) Steel globules and dendrites embedded in a copper matrix; (b) steel matrix containing a globuleof copper.

Table 6Quantitative evaluation of the percentage of both ferrous- and copper-rich phases in the FZ of joint 2: Cu–AISI 304 welding

Zone Copper side Centre Steel side Total

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

Zone 1 81 19 83 17 31 69 65 35Zone 2 78 22 88 12 47 53 71 29Zone 3 75 25 88 12 55 45 72.6 27.4Zone 4 53 47 67 33 41 59 53.6 46.4Zone 5 85 15 34 66 23 77 47.3 52.7

Fig. 11. Microhardness profiles measured at three different depths of the beadof joint 2: Cu–AISI 304.

copper base material was estimated to 110–120 HV0,1 comparedto 220 HV0.1 of the AISI 304 stainless steel.

3.3. Joint 3: CuCrZr–AISI 316L

The third sample was a double-pass welded joint obtained bywelding a plate of copper alloy C18100 (Table 1) to one of typeAISI 316L stainless steel. The geometrical characteristics of theFZ are described in Table 7, where numbers 1 and 2 correspondto the first and second pass carried out on opposite sides of thesample.

Fig. 12 shows a schematic profile and two micrographs ofFZ–base material interfaces.

Fig. 13 shows a microcrack at the FZ–steel interface intowhich melted copper has penetrated.

I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173 169

Fig. 12. Weld bead of joint 3: Cu alloy–AISI 316L. Left: FZ–copper alloy and right: FZ–AISI 316L interfaces.

Table 7Geometrical parameters of the FZ of joint 3: Cu alloy–AISI 316L welding

Maximum depth 1 (mm) 45.6Maximum depth 2 (mm) 31.7Maximum width 1 (mm) 12.7Maximum width 2 (mm) 10.8Depth to width ratio 1 3.6Depth to width ratio 2 2.9

Fig. 13. Microcrack at the interface between the FZ and steel of joint 3: Cualloy–AISI 316L.

The micrograph of Fig. 14a shows the microstructure of thebead at the intersection between the two fusion zones, whileFig. 14b shows a micrograph of the bottom zone of the joint. Nodendritic structure was identified in this seam.

The results of the area fraction evaluation of the ferrous-and copper-rich matrices are summarized in Table 8. The imageanalysis was carried out in six positions of the bead at a mag-nification of 100×. The average area fraction of the copper-richphase in the FZ is 51%.

Elemental point analyses and concentration profiles forFe(K), Cr(K), Ni(K), Si(K), Mo(K) and Cu(K) were carried out.Fig. 15 shows a SEM micrograph of the FZ–steel interface (tothe left), and the two optical micrographs (to the right) whichhighlight the details; the presence of porosities and of nonho-mogeneous structures can be noted.

The point analyses shown in Fig. 15 are related to the compo-sition of the zones A, B and C marked in the SEM micrograph.It is important to note that the globule showing light contrast

Table 8Quantitative evaluation of the percentage of the ferrous- and copper-rich phases in FZ of joint 3

Zone Copper side Centre Steel side Total

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

% Cu richphase

% Fe–Cr–Nirich phase

Zone 1 62 32 75 25 14 86 50.3 49.7Zone 2 67 33 75 25Zone 3 66 34 72 28Zone 4 45 55 76 24Zone 5 60 40 68 32Zone 6 18 82 75 25

Fig. 14. Micrographs of the FZ in joint 3: Cu a

47 53 63 3724 76 54 4630 70 50.3 49.719 81 49 5119 81 37.3 62.7

lloy–AISI 316L. (a) Centre; (b) bottom.

170 I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173

Fig. 15. Micrographs and point intensity profiles of the FZ–steel interface of joint 3: Cu alloy–AISI 316L. Left: SEM micrograph; centre: optical micrograph; right:detail.

(B) contains darker small features. Zone A shows a composi-tion similar to that of 316L stainless steel, with some traces ofcopper. The darker features inside the globule (B) shows higherconcentrations of copper and a lower presence of the other ele-ments. The C zone is rich both in the elements from the 316L

Fig. 16. Microhardness profiles measured on three different depths of the beadof joint 3: Cu alloy–AISI 316L.

alloy and in copper. The EDS analysis of the material insidethe crack shows the presence of copper. Microhardness profilesmeasured on three different depths of the fusion zone are shownin Fig. 16. The microhardness mean value of copper base mate-rial was in the range 120–130 HV0,1, while 280 HV0.1 was theaverage value of the AISI 316L stainless steel. Also in this thirdcase, measures confirm the heterogeneity of the microstructure:the hardness peaks in the FZ near to the steel side correspond toiron rich phases, while, on the opposite side, the lower micro-hardness values are due to a higher amount of copper. From amacroscopic point of view, the microhardness profile confirmsthe macrosegregation phenomena involving Fe and Cu and cov-ering the whole FZ.

4. Discussion

4.1. Phases present in the welds

The binary Fe–Cu equilibrium diagram, Fig. 17 [8], showslimited solid solubility of copper in austenite and of iron in cop-per. There are no intermetallic phases in the system. Even though304 and 316 stainless steels contain high amounts of chromium

I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173 171

Fig. 17. Fe–Cu phase diagram [9].

and nickel one can assume, as a first approximation, that this doesnot extend the solid solubility of copper in the austenite matrix.Only two phases, one copper rich and one iron rich have beenfound in the microstructures. Furthermore, it seems, from theEDS analyses, for example in Fig. 6, that no significant redistri-bution of chromium and nickel has occurred during the welding.The reason is, most likely, that the time has been too short forefficient mixing of the steel and copper melts which means thatiron rich melt areas have also been rich in chromium and nickel.

4.2. Melting during welding

The area fraction analysis of the fusion zone have shownvalues ranging from 51 to 62%. Given the limited solubilities ofcopper in austenite and iron in copper and the similar densities ofcopper and austenite, this shows that the amount of copper andaustenite that has melted during the welding is similar despitethe difference in melting point of almost 500 K. This is typi-cal for high energy welding processes and in accordance withobservations of Wei et al. [6].

4.3. Solidification microstructures

The microstructures of all the fusion zones are characterizedby copper-rich areas with primarily globular austenite crystalsascstcaltp

of directional solidification from the plates except for joint 3,Fig. 12. This nucleation seems to have been very copious andcan be attributed to the high cooling rate and corresponding highlocal undercooling. Due to the steep concentration gradients, theiron globules have normally not been able to grow long enoughto develop into dendrites. In some iron rich areas, large austenitegrains have formed and in these areas a fine segregation patternof copper can be found as can be seen in Figs. 2b, 4b (zone D),10b, and 14a. The iron rich areas also contained larger areasof entrapped copper melt in which small globular and dendriticaustenite grains have solidified as can be seen in Fig. 10b. Theseentrapped copper areas, that were the last ones to solidify, some-times contained shrinkage pores as seen in Fig. 9b.

4.4. Heat affected zones

The heat-affected zone was negligible in the first samplewhile it was few tens of �m in the second and third specimen. Inthese two cases the HAZ was thicker at the central zone of thebead where the two welding passes overlapped. These observa-

Fg

nd austenite areas with globular areas of copper (Fig. 18). Thetructures are highly non-uniform showing that mixing of theopper and austenite melts have not been achieved during thehort fusion period. There are, however, clear evidence of strongurbulence; Figs. 6 and 15 show that there is not a gradualontinuous concentration variation between the copper andustenite plates but large local concentration variations. Thesearge local concentration gradients and high cooling rate has con-rolled the solidification microstructure. Austenite grains havereferentially nucleated in the melt and there is little evidence

ig. 18. Schematic representation of the two main different morphologies oflobules observed.

172 I. Magnabosco et al. / Materials Science and Engineering A 424 (2006) 163–173

Table 9Summary of different characteristics of the three joint analysed

Power input(kW)

Thickness(mm)

Number ofwelding pass

Heat affectedzone (�m)

Microstructure Defects

Joint 1 13.5 33–29 1 – Fine dendrites, globules Copper penetrations in tograin boundaries of thesteel BM

Joint 2 18 50 2 <10 Extrafine dendrites globules Cracks, shrinkage,porosities

Joint 3 19.5 71 2 <10 No dendrites, fine globules Cracks, shrinkage,porosities

tions are in agreement with the different power inputs used inthe three cases and with the heating effect induced by the secondwelding pass.

Except for some copper penetration into the grain boundariesof the unmelted steel, no significant defects were found in thefirst joint which is consistent with the results of Tosto et al. [7]where similar welding parameters and thickness of plates wereused.

In the second and third joints, copper diffusion resulted inembrittlement of the austenitic grain boundaries; in some casesthis phenomenon promoted microcracks induced by thermalstresses (Figs. 9a, 12, 13, and 15). Most of these microcracksshowed traces of copper and it can be assumed that liquid cop-per has penetrated these boundaries but solidification shrinkageprevented efficient filling of these microfissures.

4.5. Defects, geometry and process parameters

By varying process parameters, chemical composition andthickness of the plates, different joint characteristics (summa-rized in Table 9) can be obtained. The absence of defectssuch as porosity and microfissures in joint 1 leads (from amicrostructural point of view) to sound weld. Copper pene-trations seems do not affect the static rupture: in similar con-ditions (absence of porosity and microcrack, presence of cop-pstatcmm

mdacoatmitp

5. Conclusions

Microstructural characterisation of the EBW process ofthree different dissimilar welded joints was carried out. Thesamples, characterized by different thicknesses, were obtainedby welding plates of copper alloys to three different plates ofaustenitic stainless steel using different process parameters.All the samples showed a microstructure characterized by amixture of two non-equilibrium phases, one rich in Cu andthe other one austenite with Fe, Cr and Ni. This was attributedto different factors such as low solubility in the solid statebetween iron and copper, high cooling rate typical of thiswelding technique and strong convection occurring in thekeyhole. Due to the different melting temperatures of thewelded alloys and the steep concentration gradients arisingfrom the strong convection, the microstructure detected showedsteel globules and dendrites embedded in a copper matrix andareas of austenite matrix containing globules of copper. Whilethe first sample did not reveal defects, except for some copperpenetrations into the grain boundaries of the unmelted steel, thesecond and the third joint showed porosity in the copper-richphase and intergranular microfissures at the interface betweenthe fusion zone and the stainless steel. The porosity was dueto the shrinkage of copper trapped in the solidified steel; themicrofissures were caused both by the embrittlement inducedby the copper grain boundary wetting and thermal stresses.

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R

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er penetrations in the steel grain boundary), Tosto et al. [7]howed a tensile strength very close to the copper value. Onhe other hand, it can be reasonably supposed that porositynd microfissures in joints 2 and 3 may reduce the ductility ofhe weld. Furthermore, copper penetrations and microfissuresould reduce the fatigue strength. In particular, thermal fatigueay be very critical because of the thermal instability of theicrostructure.The difference in chemical composition causes only a

icrostructure refinement: the joint 3 is characterized by mostifferences in base materials composition and it showed anbsence of dendrites and a reduction of globule dimensionompared with joints 1 and 2. Since joints 1 and 2 were carriedut using the same materials, defects in joints 2 and 3 arettributed to the process and geometry parameters; a highhickness promotes, for example, a higher cooling rate that

ay prevent the full filling of microfissures by copper; anncrement of power input increases the size of FZ and modifieshe convective motion which may affect the presence oforosity.

The investigation demonstrates that dissimilar metal weldingan be very critical. The differences observed between the firstoint and the latter two joints show that sound welds can bebtained only if an accurate optimisation of the process parame-ers (i.e., power, number of passes, speed, etc.) are performedefore the final welding operation. Further experiments haveo be carried out in order to evaluate the static and dynamic

echanical properties of these joints.

cknowledgements

The authors gratefully acknowledge the support provided byrof. Alberto Tiziani (DTG, University of Padova), Dr. E. Dellaovere and Mr. G. Mazzacavallo (DTG, University of Padova),r. G. Corniani and Mr. M. Festa (Zanon factory, Schio (Vi)).

eferences

1] P.S. Wei, F.K. Chung, Metall. Mater. Trans. B 31 (2000) 1387–1397.

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[2] Z. Sun, R. Karppi, J. Mater. Process. Technol. 59 (1996) 257–267.[3] E. Zumelzu, C. Cabezas, J. Mater. Process. Technol. 57 (1996) 249–252.[4] T.A. Mai, A.C. Spowage, Mater. Sci. Eng. A 374 (2004) 224–233.[5] P. Bala Srinivasan, V. Muthupandi, W. Dietzel, V. Sivan, Mater. Des. 27

(2006) 182–191.[6] P.S. Wei, Y.K. Kuo, J.S. Ku, Trans. ASME 122 (2000) 626–631.

[7] S. Tosto, F. Nenci, G. Hu Jiandong, F. Corniani, Pierdominici, Mater. Sci.Technol. 19 (2003) 519–522.

[8] ASM HANDBOOK, 1991. Properties and Selection: Nonferrous Alloysand Special-Purpose Materials, vol. 2.

[9] http://www.matweb.com.