Materials Science and Engineering A 424 (2006) 117–127

Microstructural, compositional and residual stress evaluation of CO2

laser welded superaustenitic AISI 904L stainless steel

A. Zambon a,∗, P. Ferro b, F. Bonollo b

a DIMEG, University of Padova, Via Marzolo, 9 I-35131 Padova, Italyb Department of Management and Engineering, University of Padova, Str.lla S. Nicola, 3 I-36100 Vicenza, Italy

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

Abstract

CO2 laser welding was performed on AISI 904L superaustenitic stainless steel sheets, with optimised processing parameters determined bymeans of melt run trial evaluations.

X-ray diffraction phase identification and light microscopy confirmed that the weld structure is fully austenitic and dendritic. A hardness increasein the weld bead with respect to the parent metal occurred and was related to both the microstructural refinement induced by a rapid cooling of thefusion zone and the presence of nanometric scale precipitates observed by TEM in the weld bead.

d

©

K

1

(san

(aa

tsodhm

0d

Residual stresses were determined by means of X-ray diffraction, exhibiting tensile stresses, close to the yield strength, in the longitudinalirection in the weld bead, while the stresses were compressive in the transverse direction and in the base material.

Tensile tests showed that welded specimens retained strength and ductility values comparable to those of the base material.2006 Elsevier B.V. All rights reserved.

eywords: Superaustenitic stainless steels; Laser welding; Residual stress; Tensile strength; AISI 904L

. Introduction

The evolution of the huge class of austenitic stainless steelsASS) started about 80 years ago and at first included theuperaustenitic subclass. The term superaustenitic relates toustenitic stainless steels containing high amounts of molybde-um, chromium, nickel, nitrogen and possibly manganese [1–4].

Table 1 summarizes and compares the typical compositionswt.%) of the widely used plain ASS AISI 304 and AISI 316nd that of the superaustenitic grade studied in the present work,long with those of a selection of other corrosion resistant steels.

The wide utilization of plain ASS, due to both their satisfac-ory corrosion resistance and their cost-effectiveness, has led totudies and researches concerning their applications in a varietyf situations and environments, so that a satisfactory amount ofata is available concerning their shaping, welding (even withigh power density facilities) and eventually post weld heat treat-ents [5].

∗ Corresponding author. Tel.: +39 049 8275502; fax: +39 049 8275500.

However, when superior corrosion resistance at moderate orhigh temperatures are required, high nickel containing alloys arefrequently adopted [6–11].

Superaustenitic stainless steels are particularly interestingbecause they bridge between relatively cheap ASS and expen-sive Nickel base superalloys, when high corrosion properties arerequired at moderately high temperatures [12].

The microstructure of superaustenitic stainless steels (SASS)consists of a fully austenitic structure in the solution-quenchedcondition. However the high amount of alloying elements, Mo,N, Cr can enhance the precipitation of intermetallic phases suchas �, R and � if soaked in the “critical” temperature interval550–800 ◦C [5,13]. Moreover, if the carbon content is higherthan 0.02% and the steel is held for sufficient time intervalsat temperatures between 450 and 800 ◦C, carbide precipitationoccurs (sensitisation effect), with negative consequences on thecorrosion resistance. When corrosion has occurred, also themechanical strength is reduced, due mainly to the well-knownphenomenon of intergranular corrosion.

In order to preserve their superior corrosion resistance, whichmay be impaired by tensile stresses (including residual stresses),

E-mail address: a.zambon@unipd.it (A. Zambon). stress relieving can be carried out at temperatures below the

921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.03.003

118 A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127

Table 1Typical chemical compositions (taken from ASTM specification for plate, sheet, and strip when available or from company data sheets) of austenitic, superausteniticand duplex stainless steels (wt.% maximum, unless range or minimum is indicated)

Alloy name UNS number Cr Ni Mo Cu N C max Mn max Si max

304 S30400 18.0–20.0 8.0–0.5 – – 0.10 max 0.08 2.00 1.00316 S31008 16.0–18.0 10.0–4.0 2.0–3.0 – 0.10 max 0.08 2.00 1.00904L N08904 19.0–21.0 24.0–6.0 4.0–5.0 1.0–2.0 <0.15 0.02 2.00 1.00254 SMOa S31254 19.5–0.5 17.5–8.5 6.0–6.5 0.50–1.00 0.18–0.22 0.02 0.45 0.33654 SMOa S32654 24.0–6.0 21.0–3.0 7.0–8.0 0.30–0.60 0.45–0.55 0.02 2.0–4.0 0.282205 S32205 22.0–3.0 4.5–6.5 3.0–3.5 0.18 0.14–0.20 0.030 1.41 0.40SAF 2507b S32750 24.0–26.0 6.0–8.0 3.0–5.0 0.50 0.24–0.32 0.030 – 0.19

a Trade name of Avesta Sheffield, AB.b Trade name of Sandvik, AB.

sensitisation temperatures or below the temperature intervals atwhich intermetallic phases precipitate.

In order to restore the corrosion resistance, the heat treatmentsequence that is carried out on such alloys is solution heatingand quenching, to be performed when a sensitisation effect mayhave occurred for any reason.

It is well known that a drawback of ordinary ASS is theirmoderate resistance to pitting corrosion, which is particularlyevident when such alloys are in contact with a halide-containingcorrosive medium. The increase in molybdenum content as wellas in nitrogen can dramatically improve the pitting resistance.This effect can be pointed out by the pitting resistance equivalent(PRE) number [5], which depends on the chromium as well ason the molybdenum and nitrogen content.

In Table 2 a comparison is given of the PRE numbers as wellas of the corresponding critical pitting temperatures (CPT). Itcan be easily noticed that both values are well improved for thesuperaustenitic steel.

It is well known that austenitic stainless steels suffer fromstress corrosion cracking (SCC), that is the development ofcracks while the alloy is in contact with a corrosive environ-ment (most often a chloride containing solution), and a tensilestress is acting on the workpiece or test specimen. Different testshave been developed to obtain comparative resistance in stresscorrosion conditions. In particular if dilute NaCl aqueous solu-tion drips onto a heated specimen whose working temperature issics

c

TMs

A

AA9

superaustenitic stainless steels is far higher than that of ordi-nary austenitic stainless ones and also than that of duplex steels.

On the other hand the proneness to intermetallic phase pre-cipitation of superaustenitic and duplex steels exhibiting similarcorrosion resistances, is lower for the superaustenitic grades, ascan be evaluated on the basis of the precipitation index (P.I. = Cr(wt.%) + 3.3Mo (wt.%)) [14], due to the overall lower amountof chromium and molybdenum.

Sheets and pipes of all the austenitic and superaustenitic stain-less steels can be welded by plasma arc welding, TIG or MIG,with suitable filler metals. However, different problems can arisewhen these materials are welded.

Lu et al. [15] found that the gas tungsten arc welding reducedthe pitting corrosion resistance in the weld metal and heataffected zone (HAZ) of an AISI 304 austenitic stainless steel; onthe contrary, laser-beam welding improved it. High heat-input(and low cooling rate) was likely to induce the segregation ofalloying elements and formation of Cr-depleted zones, resultingin degradation of the resistance to localized corrosion.

Intermetallic phases were found by McPherson in the fusionzone (FZ) of laser and submerged arc welded joints of 316LNand duplex 2205 stainless steels [16]. The authors admitted thatthe presence of those intermetallic phases (� and Cr2N) in laserwelded joints was unexpected, and could only have occurred dueto some high temperature segregation effect.

Heino et al. [17], by investigating the welding of the super-atclptw

atpatGpDh

et in order that a drop evaporates before the next hits the spec-men surface, the stress corresponding to the onset of crackingan be obtained for different alloys, which can give a relativeequence of resistance to SCC.

In [4] such stress values among different materials wereompared, and it was shown that the resistance to SCC for

able 2inimum values of PRE e CPT for selected austenitic/superaustenitic stainless

teels

lloy PREa CPT (◦C)b

ISI 304 18 3ISI 316 23 1704L 32 44

a PRE = wt.% Cr + 3.3(wt.% Mo) + 16(wt.% N).b Ref. [5].

ustenitic stainless steel Avesta 654 SMO, detected in the HAZwo different type of precipitates, � and R. The absence of pre-ipitates in the region near the plate surface, closest to the fusionine, was justified by the cooling rate which was fast enough torevent precipitate formation. However, closer to the plate cen-res, the welding process gave rise to a temperature exposure athich precipitation occurred.Ames et al. [18] compared the austenitic (316L), super

ustenitic (254 SMO) and super duplex (SAF 2507) weld proper-ies produced using GTAW flux. The microstructure, mechanicalroperties and corrosion resistance of such welds were comparedlso to autogenous welds produced using the current indus-ry standard gas (argon). In all cases, the welds produced withTAW flux exhibited equivalent or improved microstructure androperties when compared to the autogenous non-flux welds.epending on the amount of cold work present, the strength andardness in the HAZ and FZ softened due to the weld thermal

A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127 119

cycle. Microhardeness traverses across welds examined in [18]showed this softening effect.

It has been demonstrated that ASS can be easily welded byhigh power density sources, such as focussed laser or electronbeams [19–22]. The major advantages of these welding pro-cesses are a small HAZ, which cools very rapidly with verylittle distortion, and a high depth-to-with ratio for the fusionzone. Generally, laser beam welding (LBW) involves many vari-ables; laser power, welding speed, defocusing distance and typeof shielding gas, any of which may have an important effect onheat flux and fluid flow in the weld pool. This in turn will affectpenetration depth, shape and final solidification structure of thefusion zone.

Bonollo et al. [23] investigated the microstructural andfatigue properties of different superaustenitic stainless steels(AISI 904L, 254 SMO, AMANOX 1.3964) obtained by means ofplasma arc welding (PAW) and electron beam welding (EBW).The microstructures were fully dendritic and austenitic althoughthe interdendritic material was more rich in ferrite stabilizing ele-ments (Cr, Mo). The HAZ was very narrow in the electron beamwelded joints due to the high cooling rate typical of that weld-ing technology. The fatigue strength of such welded joints wasdetermined to be superior to that of joints of ferritic–pearliticsteels. However, in that work, specific information about inter-metallic phases, microhardness profiles, residual stress and staticmechanical properties was not given.

tttn

i9

2

olbbP

abtFg

Table 4Parameters and results of laser melt run trials and welding

Traverse speed(mm/s)

Heat input(kJ/m)

Completepenetration

Test 1 20 175 Yes/deepTest 2 30 117 Yes/discontinuousTest 3 40 88 NoWelded joint 25 140 Yes/deep

Focus position (relative to surface of samples), f = 0 mm, beam power,P = 3.5 kW, beam size on surface of specimens 0.35 mm.

The penetration beads were examined by radiographic obser-vations (using an XRC 600 X-ray controller machine, with accel-erating voltage of 110 kV and a current of 3 mA) and checkedfor surface discontinuities with dye penetrant. Finally they weresectioned to perform metallographic observations and check ofpenetration depths (see Table 4 for results).

Gas porosities (probably due to an excess of nitrogen in themolten pool) as well as discontinuous undercuts occurred along-side the bead, even in the case of complete penetration.

It was therefore decided to produce the weld beads at a tra-verse speed of 25 mm/s, moreover adopting helium instead ofnitrogen as the shielding and plasma removing gas. The relevantparameters for welding are summarized in Table 4.

Again prior to further examination steps, the above describednon-destructive tests were performed, which confirmed in thiscase the absence of any detectable porosity and undercuts. Inparticular, following the standard EN 462/1, defects with adimension greater than 150 �m were not detected.

The butt-welded joints were thus investigated by means oflight microscopy, microhardness measurements, TEM observa-tions, residual stresses determinations and tensile strength tests(after machining of the nail head and of the root of the weld beadfollowed by polishing).

The etchant composition for light microscopy is given inTable 5.

Microhardness (DPH-HV0.2) tests were performed on a trans-vpa

TE

E

TC stand

A

AA .0

It can be pointed out that only few researchers have inves-igated in the past the effects of high power density weldingechniques on superaustenitic stainless steels. Specific informa-ion on the high-energy beam welding of AISI 904L alloy haveot yet been fully established.

The present work is focussed on the microstructural, mechan-cal and residual stresses assessment of laser beam welded AISI04L (UNS N08904) steel.

. Material and methods

Three-millimetre thick sheets of AISI 904L, the compositionf which is summarized in Table 3, were machined with paral-el edges in order to afford a narrow transverse gap (<0.1 mm)etween the components to be welded. The sheets were weldedy the “SITEC” laboratory in Lecco (Italy), belonging to theolytechnic of Milan.

Prior to the welding, a series of melt run trials were performedt constant power density focussed on the sheet surfaces. Theeam power of a 6 kW CO2 laser was set at 3.5 kW for all theests, while the traverse speeds ranged between 20 and 40 mm/s.or all penetration tests nitrogen was adopted as the shieldingas.

able 3hemical composition of the AISI 904L steel used in the tests compared to the

lloy C Mn Cr Ni

ISI 904L 0.013 1.45 19.52 25.03ISI 904L (standard) <0.02 <2.0 19.0–21.0 24.0–26

erse section of the weld bead (Fig. 1(a)), along three directionsarallel to the surface of the sheet (near the head, in the centrend near the root of the bead) in order to identify the possible

able 5tchant used for light microscopy examinations (soaking time: 60 s)

tchant “Beraha”100 ml H2O20 ml HCl0.5–1.0 g K2SO4

ard one

Mo N Cu Si P S

4.27 0.05 1.41 0.23 0.025 0.0014.0–5.0 <0.15 1.0–2.0 <1.00 <0.045 <0.035

120 A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127

Fig. 1. Scheme of the microhardness tests positions (a) and investigated zonesfor residual stresses measurements (b) (specimen dimensions: 80 mm × 30 mm).

effects of microstructural heterogeneities both in the bead, andin the parent metal.

Thin foil specimens were obtained by sectioning perpen-dicularly to the bead, gently grinding with emery papers to athickness of ∼0.07 mm, ending with final polishing on a nap witha 1 �m diamond suspension, punching out a 3 mm diameter diskand final etching in a dual-jet thinner with a HClO4-10 vol.% inethylene glycol solution, refrigerated at −10 ◦C.

The residual stress determinations were carried out on thesurface of the welded sample, on top of the bead, in the HAZ,in the base metal 7 and 12 mm aside the bead central line.

Such determinations were performed by X-ray diffraction,according to the well known “sin2 ψ” method, using Cr K�radiation, diffracted by the {2 2 0} crystallographic planes ofthe lattice of Austenite. Diffracted intensities were acquiredfor tilting angles of 0◦, 15◦, 30◦, 45◦ and 60◦ in four posi-tions, as shown in Fig. 1(b). The spot size, whose dimensionswere 1.5 mm × 1.5 mm, was obtained by masking the area sur-rounding the intended spot with paint containing TiO2. Theinvestigated positions were at the bead centreline, at bead border(“intermediate zone”), at 7 and 12 mm from the bead centreline.

Finally, mechanical tests were carried out on short specimensmachined according to the ASME IX standard, on a 100 kN MTStest machine.

3. Results and discussion

3

sw

residual stress determinations were performed and from whichspecimens for metallographic observations and microhardnessmeasurements were obtained. Fig. 2(d) shows also the ImageQuality Indicator (W 10 FE) used in the radiographic observa-tions in order to evaluate the minimum detectable defect size.It is clearly noticeable the absence of any detectable poros-ity (diameter greater then 150 �m) or undercut, especially inthe central region of the samples. The sides (laser-in side andlaser-out side of the weld beads) were subsequently removed toperform mechanical tests (Fig. 3).

3.2. Macrostructure

The macrostructure of the weld and proximal parent metal isvisible in Fig. 4. A slight depression of the head of the bead isnoticeable, probably due to insufficient matching of the two fac-ing edges. The vertical matching appears good; this contributesto the symmetry of the bead. Due to the small thickness of thesheets and to the position of the focus of the beam, the depth-to-width ratio of the fusion zone (FZ) is relatively small (withreference to Fig. 4 the depth-to-width ratio is 1.36).

3.3. Microstructure

The micrographic examination shows, first of all, that theextension of the heat affected zone is negligible (Fig. 5(a)). Thisiici

cddrd

b(afiw

f

C

N

iNsoIbsAr

.1. Radiographic inspection

Fig. 2(a) and (b) shows the radiographies of the weldedamples from which the tensile test specimens were obtained,hile Fig. 2(c) shows that of the welded sample on which

s the typical positive effect associated to a highly focussed weld-ng heat source, such as a laser beam, producing extremely highooling rates [22,24]. The extension of the heat affected zone isn the order of 10 �m.

In the weld bead, a fully dendritic microstructure (Fig. 5(b))an be detected, with well-developed primary arms and clearlyistinguishable short secondary arms. From image analysiseterminations, it has been seen that the darker (dendritic) phaseepresents 62% of the bead area, and the lighter one (interden-ritic) the remaining 38%.

On the other hand, when the weld bead was investigatedy means of X-ray diffraction (XRD) for phase identificationFig. 6), with a tilting and rotating sample holder in order tovoid or at least minimize any texture effect due to the solidi-cation regime, just the reflections corresponding to Austeniteere detected.In fact, from the composition given in Table 3 and from the

ollowing formulae [15]:

req = %Cr + %Mo + 0.7 %Nb

ieq = %Ni + 35 %C + 20 %N + 0.25 %Cu

t can be calculated that the steel under investigation has a Creq toieq ratio equal to 0.88. This means, as shown in Fig. 7, that the

olidification, under equilibrium conditions (but also in presencef high cooling rates), takes place in an austenitic mode [22,26].n literature [23,26,27], the apparent presence of two phases haseen detected, and has been attributed to the fact that the firstolidifying grains (i.e. the dendrites) are certainly enriched inustenite stabilizing elements, while the ferrite stabilisers are

ejected to the interdendritic liquid. However, due to the very low

A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127 121

Fig. 2. Radiographies of samples 1 (a) and 2 (b) to be subsequently used in tensile strength tests; radiography of the welded plate for residual stress measurementsand subsequent metallographic observations (c); Image Quality Indicator IQI EN 462-W 10 FE (d).

Fig. 3. Sketch of the tensile test specimen.

Creq/Nieq value of this kind of steels, the interdendritic liquidalso solidifies as Austenite, with a composition only slightlydifferent from that of the dendrites.

This behaviour, in the case studied, is confirmed by EDSinvestigations, the results of which are shown in Fig. 8. The com-position, as shown by EDS spectra, is quite similar in dendriticand interdendritic regions. Some differences can be noticed insome elements (Cr, Mo and Ni) which seem to be more concen-trated in the interdendritic region.

Moreover the microstructure examination allowed evaluationof the secondary dendrite arm spacing (SDAS), whose averagevalue is given in Table 6, and which can be correlated with thelocal cooling rate, depending on the heat input (i.e. the ratio

122 A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127

Fig. 4. Macrograph of the laser welded sample.

Table 6SDAS and related cooling rate

Material 904L/LBWSDAS (�m) 2.902Cooling rate (K/s) 6902 ± 277

between beam power and traverse speed) and on the thermaland geometrical characteristics of the sample.

According to the formula [28]:

SDAS = 64T−0.35 (�m),

T, i.e. the cooling rate (K/s) during solidification, can be com-puted achieving the result given in Table 6.

Thus, the microstructural features detected in the weld beadcan be understood only taking into account, at the same time,

the role of composition and of cooling rates induced by the laserwelding process.

(1) From the composition point of view, it is clear that both inequilibrium and under high cooling rates the alloy solidifiesfollowing an austenitic mode.

(2) The cooling rate achieved (nearly 7000 K/s) stronglyinhibits diffusion phenomena and consequently affects thepartitioning of elements. Overlapping the EDS spectra ofdendritic and interdendritic areas, a slight segregation ofsome alloy elements (Cr, Mo and Ni) is noticeable (Fig. 8)into the interdendritic regions.

(3) The relatively high content of interdendritic Austenite canalso be attributed to the high cooling rates, which makedifficult a relevant coarsening of primary Austenite den-drites (the same trend can be seen comparing laser and TIGcladding treatments [26–27]).

(4) Literature mentions that N can inhibit the formation ofsecondary dendrite arms (the threshold has been fixed at0.05%, for TIG cladding processes [27]); their presence inthe detected specimens (and also on 904L laser claddingprocesses [26]) is, again, the result of the combined actionsof composition and cooling rates (the relatively high amountof interdendritic liquid perhaps makes easier the develop-ment of such arms).

3

dzje(tscdc

d bea

Fig. 5. Micrograph of the interface between wel

.4. TEM examinations

The compositional heterogeneity of dendritic and interden-ritic areas probably affected the corrosion rate of the twoones during the electrochemical jet thinning. As a consequenceust restricted areas of the weld bead were transparent to thelectron beam. Very small precipitates of the nanometric scale20–40 nm) can be observed in Fig. 9(a), along with disloca-ion loops, some of which are pinned by such precipitates. Theelected area diffraction pattern (Fig. 9(b)), with spotted rings,onfirms that the region is crystalline, with small crystallineomains. This substructure is obviously due to the effect of theomposition of the region and is presumably deeply affected by

d and parent metal (a) and of the weld bead (b).

A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127 123

Fig. 6. Diffraction spectrum of the cross-section of the weld bead with identifi-cation of the diffraction reflections.

the local cooling rates. In the light of such parameters it can-not be established if such nanometric features originated duringsolidification or in the subsequent cooling.

3.5. Microhardness tests

As can be seen in Fig. 10 the microhardness values (HV0.2)increase moving from the base material across the weld beadtowards its opposite side. During the solidification of the FZ, thematerial generally loses its original strength induced by strainhardening. On the contrary, the formation of bainitic or marten-sitic phase in the FZ in the case of carbon or low alloy steels

Fig. 7. Pseudo-binary diagram for the prediction of equilibrium microstructurein stainless steels [25].

promotes an increase of hardness. In single-phase materials, thishardness increment in the FZ may be induced by precipitationhardening effects, residual stresses or microstructure refine-ment due to the rapid cooling of the weld pool. Microhardness

Fig. 8. Results of EDS analysis carried out in the dendritic (

grey spectrum) and interdendritic (black line) regions.

124 A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127

Fig. 9. (a) TEM micrograph of thin foil specimen obtained from a transverse section of the weld bead showing the distribution of nanometric-scale precipitates.Dislocation loops are clearly distinguishable in the matrix; (b) corresponding selected area diffraction pattern showing spots arranged in rings demonstrating thenano-sized structure of the investigated area.

Fig. 10. Microhardness profiles of the weld bead and parent metal.

measurements on as-welded and stress-relieved (380 ◦C, 1 h)specimens (Fig. 11) showed the same trend so that residualstresses cannot explain the observed increment of hardness inthe pool. On the other hand, microhardness profiles in weldedjoints obtained with different welding speeds show an incre-ment of strength in the FZ induced by a finer microstructure(Fig. 12). On the basis of these results and of TEM examinations,it is likely that the hardness increment in the FZ is a conse-quence of the presence of both the observed precipitates and

Fig. 11. Comparison between the microhardness profiles near the head of thebead of as-welded and stress-relieved specimens.

grain boundaries. They interact with (and hinder) the movementof dislocations necessary for the deformation of the materialunderneath the indenter while the diamond pyramid hardness(DPH) test is being carried out.

3.6. Residual stresses

When first and second-order residual stresses act in a crys-talline workpiece, it is possible to evaluate the shifting of theintensity peaks of the diffracted signal deriving from a beam ofX-rays impinging on the surface of the polycrystalline sample,when the sample is tilted with respect to the original orienta-tion. Though in principle any original tilting angle ψ could beadopted, most frequently it is taken equal to zero. Referring toFig. 13(a) an evident shifting of the centroids of the diffractedintensity peaks towards increasing values of the abscissa θ isnoticeable. As θ is related to the interplanar distances of thediffracting planes through the Bragg’s law, the relative positionshown in Fig. 13(a) of the aforesaid centroids shows a decreaseof the interplanar distances of crystals oriented at higher tiltingangles to the original ψ = 0 direction. This observation can beeasily related, under the hypothesis of an elastically behavingpolycrystalline isotropic material (and adopting average elasticconstants), to the mean stress acting parallel to the surface, in a

Fh

ig. 12. Influence of welding speed on microhardness profiles obtained near theead of the bead.

A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127 125

Fig. 13. Scheme of the “sin2 ψ” method: intensity vs. 2θ profiles (a), “ε” vs.“sin2 ψ” plot (b).

Table 7Residual stresses in the weld bead and in the parent metal (spot size:1.5 mm × 1.5 mm)

Longitudinal stress (MPa) Transverse stress (MPa)

Weld bead 239 ± 10 −345 ± 9Intermediate zone 128 ± 16 −113 ± 57 mm −364 ± 22 −149 ± 912 mm −417 ± 18 −181 ± 14

direction perpendicular to the axis of tilting, within the layer con-tributing to the diffracted intensity. In the case shown in Fig. 13the aforesaid stress component is compressive. An example ofthe “ε” (that is (d(�2 2 0)ψ − d(�2 2 0)0)/d(�2 2 0)0) versus “sin2 ψ”plot is shown in Fig. 13(b).

The residual stresses values are summarized in Table 7. Asevident, the longitudinal stress on the surface of the head of theweld bead appears to be close to the yield stress (YS = 270 MPa,hot rolled plate [4]; YS = 315 MPa, measured value in the as

Table 8Results of the tensile test (ultimate tensile strength (UTS) and strain-to-failure)

Measured values Nominal and measured valuesof base metal

UTS (MPa) Strain-to-failure (%)

UTS (MPa) Strain-to-failure (%)

Sample 1 613 >60 >550 >50Sample 2 619 >60 584 (measured)

received material), neutral in the transition zone with the parentmetal and compressive in the base material.

The measured residual stresses in the weld bead and in theintermediate zone are in agreement with the typical behaviourobserved and calculated in butt-welded joints of different mate-rials which do not present phase transformations during welding(i.e., aluminium or nickel alloys) [29,30]. In fact, the vol-ume change induced by phase transformation (such as themartensitic transformation which affects carbon steels) andthe transformation plasticity phenomenon [31–33] are likely torelieve high tensile stresses in the FZ of multi-phase materials[34–36].

The high compressive residual stresses measured on the sur-face of the parent material are probably due to the rolling effectspresent in the as-received sheet metals.

3.7. Tensile strength

First of all, it must be pointed out that the rupture of boththe specimens occurred in the base metal, far away from theweld bead and the intermediate zone (Fig. 14). Therefore, thetensile strength of the weld bead exceeds that of the base metal(Table 8). This is not surprising, as it is known that hardness andtensile strength values are related.

Though the ratios of the specimens dimensions did not followtts

ssdd

j

Fig. 14. Tensile spe

he standard indications, it can be pointed out that the deforma-ion to rupture is likely to have been retained in the weldedample as a whole (Fig. 15).

A keen observation of the surface of the test specimen clearlyhows that the weld bead, which before the test was absolutelymooth, shows a sequence of sags (Fig. 14) due to differentialeformations occurred in the weld bead, as a consequence ofifferent orientations of the microstructural constituents.

Since both the strength and deformation properties of theoint as a whole are very close to those of the base metal, it can

cimen No 2.

126 A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127

Fig. 15. Stress–strain curve for specimen 2.

be concluded that the joint retained a toughness comparable tothat of the base material.

4. Conclusions

Welding parameters concerning power density, traverse speedand therefore heat input, as well as the shielding gas were prop-erly chosen to produce sound joints on 3-mm thick AISI 904Lsuperaustenitic stainless steel.

Non-destructive testing carried out by radiographic examina-tions and dye-penetrant confirmed the integrity of both the weldbead and the intermediate zone close to the bead.

Light microscopy examinations allowed assessment of theintegrity of the structure and to evaluate the cooling rates duringsolidification on the basis of SDAS measurements.

The weld bead structure appears to be characterised by afully dendritic microstructure, resulting from the undercoolingdue both to the rapid cooling which characterises the high powerdensity welding process and the presence of alloy elements (con-stitutional undercooling). An imperceptible heat affected zonewas found, therefore allowing the retention of high toughness.

Microhardness values in the weld bead showed an increasewith respect to those measured in the parent metal. Such anincrease has been related both to the microstructural refinementinduced by the rapid cooling in the FZ and to nano-sized pre-cipitates observed by means of TEM microscopy of thin foils

ot

tmn

la

A

e

saldature ad alta produttivita su acciai per impieghi strutturali:Analisi metallografica, di composizione e tensionale in giuntisaldati di acciai strutturali” is fully acknowledged.

The authors are grateful to Barbara Previtali, Ph.D., of SITECLecco (Laboratorio per le Applicazioni Laser, Politecnico diMilano) for performing laser welding and to Dr. AlessandroMilan, Dr. Enrico Della Rovere and Mr. Giacomo Mazzacavallofor helping in microstructural examinations.

References

[1] M. Liljas, Scand. J. Metall. (1997) 32–58.[2] G. Di Caprio, Gli acciai inossidabili, Hoepli, Milano, 1997.[3] L. Colombier, J. Hochmann, Stainless and Heat Resisting Steels, St.

Martin’s Press, New York, 1968.[4] M. Liljas, Acom 2 (1995) 1-7, Avesta Sheffileld AB (publ) (ISSN

1101-0681), http://www.avestapolarit.com/upload/documents/technical/acom/acom95 2.pdf.

[5] B. Leffler, Stainless steels and their properties, AvestaPolarit AB,Technical bulletin, www.avestapolarit.com, http://www.outokumpu.com/files/Group/HR/Documents/STAINLESS20.pdf.

[6] J. Olsson, Acom 1 (1995) 1–9, Avesta Sheffileld AB (Publ) (ISSN1101-0681) http://www.avestapolarit.com/upload/documents/technical/acom/acom95 1.pdf.

[7] C. Wolfe, E. Arnvig, W. Wasielewska, Acom 2 (1997) 1–7, AvestaSheffileld AB (Publ) (ISSN 1101-0681) http://www.avestapolarit.com/upload/documents/technical/acom/acom97 2.pdf.

[8] B. Wallen, E. Alfonsson, Acom 1 (1996) 10–14, Avesta Shef-

[

[

[

[

[

[

[

[

[

[[[[[

[

[

[

[

pecimens.The aforesaid increase in hardness agrees with the occurrence

f rupture in the base metal, which took place in all the tensileests carried out.

In any case tensile strength and deformation before rup-ure, although the tests were carried out on “short” specimens

achined according to ASME IX, appear to comply with theominal recommended values for the un-welded base material.

Therefore laser welding of superaustenitic AISI 904L stain-ess steel is fully satisfactory in terms of both microstructuralnd overall mechanical reliability.

cknowledgments

The support of the Italian Ministry for University and Sci-ntific Research under contract “Metallurgia e modellazione di

fileld AB (Publ) (ISSN 1101-0681) http://www.avestapolarit.com/upload/documents/technical/acom/acom96 1.pdf.

[9] J. Olsson, W. Wasielewska, Acom 2 (1998) 1–8, Avesta Shef-fileld AB (Publ) (ISSN 1101-0681) http://www.avestapolarit.com/upload/documents/technical/acom/acom98 2.pdf.

10] B. Wallen, A. Bergqvist, Acom 4 (1997) 1–7, Avesta SheffileldAB (Publ) (ISSN 1101-0681) http://www.avestapolarit.com/upload/documents/technical/acom/acom97 4.pdf.

11] S. Ylasaari, O. Forsen, J. Aromaa, J. Virtanen, Mater. Sci. Forum 247(1997) 83–92.

12] B. Wallen, M. Liljas, P. Stenvall, Proceedings of the Conference onApplications of Stainless Steel, Stockolm, 1992, pp. 9–11.

13] ASM Metals Handbook, Welding, Brazing and Soldering, vol. 6, 9thed., ASM Metals Park, Ohio, 1991.

14] H. Hokamoto, Proceedings of the Application of Stainless Steels’92,Stockolm, Sweden, June 9–11, 1992, pp. 360–369.

15] B.T. Lu, Z.K. Chen, J.L. Luo, B.M. Patchett, Z.H. Xu, Electrochim.Acta 50 (2005) 1391–1403.

16] N.A. McPherson, K. Chi, T.N. Baker, J. Mater. Process. Technol. 134(2003) 174–179.

17] S. Heino, E.M. Knutson-Wedel, B. Karlsson, Mater. Sci. Technol. 15(1999) 101–108.

18] N. Ames, M. Ramberg, M. Johnson, T. Johns, Stainless Steel World©,KCI Publishing BV, 2001.

19] T. Takalo, N. Suutala, T. Moisio, Metall. Trans. 10A (1979) 1173–1181.20] T. Takalo, N. Suutala, T. Moisio, Metall. Trans. 11A (1980) 717–725.21] S. Kou, Y. Le, Metall. Trans. 13A (1983) 23–31.22] A. Zambon, F. Bonollo, Mater. Sci. Eng. A 178 (1994) 203–207.23] F. Bonollo, A. Tiziani, R. Tovo, L.M. Volpone, Weld. Int. 18 (2005)

24–30.24] F. Bonollo, A. Tiziani, A. Zambon, Mater. Sci. Technol. 9 (1993)

1137–1144.25] J. Charles, Proceedings of the Conference on Application of Stainless

Steels’92, Stockholm, Jerncontoret, 1992, pp. 1108–1121.26] R. Li, M.G.S. Ferreira, M.A. Anjos, R. Vilar, Surf. Coat. Technol. 88

(1996) 90–95.27] V.S. Raja, S.K. Varshney, R. Raman, S.D. Kulkarni, Corros. Sci. 40

(1998) 1609–1625.

A. Zambon et al. / Materials Science and Engineering A 424 (2006) 117–127 127

[28] L.F. Shiau, W.G. Lo, A.W. Cramb, Iron Steelmaker 18 (1991) 57–62.[29] P. Ferro, A. Zambon, F. Bonollo, Mater. Sci. Eng. A 392 (2005) 94–105.[30] J. Canas, R. Picon, F. Parıs, A. Blazquez, J.C. Marın, Comput. Struct.

58 (1996) 59–69.[31] J.B. Leblond, G. Mottet, J.C. Devaux, J. Mech. Phys. Solids 34 (4)

(1986) 395–409.[32] J.B. Leblond, G. Mottet, J.C. Devaux, Int. J. Plasticity 5 (4) (1989)

551–572.

[33] J.B. Leblond, G. Mottet, J.C. Devaux, J. Mech. Phys. Solids 34 (4)(1986) 411–432.

[34] M. Toyoda, M. Mochizuki, Sci. Technol. Adv. Mater. 5 (2004) 255–266.

[35] M. Mochizuki, M. Toyoda, J. Phys. IV 120 (2004) 635–648.[36] Ph. Mabelly, Ph. Bourges, G. Pont, Tech. Rep. Marine Struct. 14 (2001)

553–567.