Surface residual stresses in multipass weldsproduced using low transformationtemperature filler alloys

T. I. Ramjaun*1, H. J. Stone1, L. Karlsson2, M. A. Gharghouri3, K. Dalaei4, R. J.Moat5 and H. K. D. H. Bhadeshia1

Tensile residual stresses at the surface of welded components are known to compromise fatigue

resistance through the accelerated initiation of microcracks, especially at the weld toe.

Inducement of compression in these regions is a common technique employed to enhance

fatigue performance. Transformation plasticity has been established as a viable method to

generate such compressive residual stresses in steel welds and exploits the phase transformation

in welding filler alloys that transform at low temperature to compensate for accumulated thermal

contraction strains. Neutron and X-ray diffraction have been used to determine the stress profiles

that exist across the surface of plates welded with low transformation temperature welding alloys,

with a particular focus on the stress at the weld toe. For the first time, near surface neutron

diffraction data have shown the extent of local stress variation at the critical, fusion boundary

location. Compression was evident for the three measurement orientations at the fusion

boundaries. Compressive longitudinal residual stresses and tensile transverse stresses were

measured in the weld metal.

Keywords: Transformation induced plasticity, Martensite, Residual stress, Welding, Neutron, X-ray diffraction

IntroductionThe welded joints of engineering structures are often thefeatures that limit both the service loads that can betolerated and their fatigue life. Stress at the boundarybetween the weld metal and base material caused bygeometrical changes, coupled with local microstructuralchanges, often leads to preferential crack initiation andfatigue failure at this site.1 Susceptibility to microcrackinitiation/propagation can be further exacerbated bytensile residual stresses in this region, which accumulateas a result of thermal contraction strains during coolingof the weld to ambient temperature.In order to increase the longevity of welded compo-

nents, post-weld treatments may be applied that areeither mechanical or thermal in nature. Heat treatmentsare able to relax internal stresses, while mechanicalprocesses tend to impart compression into the compo-nent surface or modify the weld toe geometry to reduce

stress concentration. As such treatments are typicallycostly or impractical, it may be preferable to modify thewelding process to minimise the occurrence of suchstresses. One method, proposed originally by Jones andAlberry,2 is to counter the thermally induced tensilestresses through exploitation of the strains associatedwith solid state phase transformation of the weld filler.Application of this mechanism has led to a number ofmartensitic welding alloys being developed.311

The transformation of austenite to martensite (cRa9)is displacive and results in a shape deformation that is aninvariant plane strain with a large shear component.12

The magnitude of the shear, in conjunction with adilatational strain, is sufficient to not only cancel thetensile stresses, but even create compression in the weldmetal. Detailed reviews of this mechanism and itsinfluence on residual stresses in welds are availableelsewhere.13,14 However, the beneficial effects of stressalleviation through transformation plasticity are depen-dant on the temperature at which the cRa9 transforma-tion occurs. If the transformation takes place above anoptimal temperature, continued thermal contraction ofthe transformed product to ambient temperature leadsto further accumulation of tensile stress. Thus, the initialbenefits associated with the transformation are elimi-nated. To avoid this problem, the weld metal should bedesigned with a low martensite start temperature (MS),typically 200uC.

1Materials Science and Metallurgy, University of Cambridge, CambridgeCB3 0FS, UK2Department of Engineering Science, University West, SE-461 86Trollhattan, Sweden3Canadian Neutron Beam Centre, Chalk River Laboratories, Chalk River,Ont. K0J 1J0, Canada4ESAB AB, Lindholmsallen 9, 40277 Gothenburg, Sweden5Materials Engineering, The Open University, Milton Keynes MK7 6AA,UK

*Corresponding author, email tir20@cam.ac.uk

2014 Institute of Materials, Minerals and MiningPublished by Maney on behalf of the InstituteReceived 20 May 2014; accepted 11 July 2014DOI 10.1179/1362171814Y.0000000234 Science and Technology of Welding and Joining 2014 VOL 19 NO 7 623

These low transformation temperature (LTT) weldingalloys have proved effective in enhancing fatigueperformance.1520 Although, judicious alloy design21,22

has been required in order to ensure that the carbonconcentration is sufficiently low to avoid the embrittle-ment associated with hard martensite. The ability ofLTT welding alloys to provide stress relief in single andmultipass welds has also been established.23 However,surface stress measurements obtained from synchrotronX-ray data are seemingly in contradiction with measure-ments made in the bulk using neutron diffraction, asthey reveal tensile stress in the weld metal surface.24,25

In this work, near surface neutron diffraction hasbeen used to probe the residual stress state in the vicinityof the free surface and these results are criticallycompared with residual stresses determined using X-ray diffraction. This study highlights the differencesbetween the residual stresses at the surface and those inthe bulk immediately below it, and serves to rationalisethe apparent disparity between the residual stress dataobtained with X-ray and neutron diffraction in previousstudies.

Experimental methodThree multipass welds were prepared from differentwelding consumables (LTT-1, LTT-2, HTT) depositedon a high strength ferritic steel plate (BP700) in order tomeasure the surface residual stresses. The specificwelding procedures are given in Table 1, where LTT-1and LTT-2 are martensitic stainless steel filler alloyswith a low MS, while HTT is a commercially availablefiller with a much higher MS.

The equation by Steven and Haynes26 was used tocalculate the MS of the welding fillers. Dilatometry wasthen used to verify the MS of LTT-1 in its undilutedstate. The measured value (16412uC) was in agreementwith the calculation. The welding alloy compositionsand predicted martensite start temperatures are shownin Table 2. LTT-2 is highly alloyed and was designed tocompensate for dilution that occurs when the filler mixeswith the baseplate.

The baseplate was prepared from 5006150615 mmsections, machined with a 60u, 8 mm deep V groovealong the long direction and a root radius of 4 mm. Theplates were clamped to the bench before mechanised gasshielded metal arc welding, which was performedhorizontally in the down hand position (Fig. 1). Allthree welding alloys were deposited using metal coredelectrode wire, with an initial preheat of 50uC and aninterpass temperature of 100125uC. The heat inputs forthe LTT and HTT welding alloys were y1?0 andy1?5 kJ mm21 respectively. Details of the weldingparameters are in Table 3. The mechanical propertiesof the welding fillers and baseplate are shown in Table 4;the LTT data are measured values.

Welds L1 and H1 were produced to compare thesurface stress distributions that develop during thecooling of a weld fabricated with an LTT filler withthose that develop using a conventional filler, with the

Table 1 Combinations of baseplate and weld ller alloys

Weld Baseplate Pass 1 Pass 2 Pass 3

L1 BP700 LTT-1 LTT-1 LTT-1L2 BP700 HTT HTT LTT-2H1 BP700 HTT HTT HTT

Table 2 Compositions (wt-%) and calculated MS of undiluted welding alloys and baseplate

Material C Si Mn Cr Ni Mo MS/uC

LTT-1 0?010?03 0?60?8 1?21?7 12?513?0 5?56?5 ,0?1 169LTT-2 ,0?02 ,1 ,2 1518 68 ,0?1 87HTT 0?12 0?65 1?50 0?70 2?80 0?85 372BP700 0?15 0?29 0?98 0?25 0?043 0?15 447

1 Photograph of weld being produced by automated gas

shielded metal arc welding

Table 3 Welding parameters

Voltage/V Current/A Shielding gas Gas flow/L min21

Welding speed/cm min21

Pass 1 Pass 2 Pass 3

24?7 y250 Arz2%CO2 18 36 30 23

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aim of correlating with fatigue data. Weld L2 wasassessed to determine whether a single capping passwould be sufficient to provide the same surface stressdistribution as a full LTT weld and hence, the desiredfatigue improvement. Surface residual stresses weremeasured with both neutrons and X-rays.

Neutron diffractionThe residual stresses measured using neutron diffractionwere performed on the L3 beamline at the CanadianNeutron Beam Centre.27 Strain scanning was performedacross a plane perpendicular to the weld at the positionalong the sample shown in Fig. 2. Measurements weremade at 0?15 and 2?5 mm below the top surface in theweld, heat affected zone (HAZ) and baseplate toascertain how the stresses vary with depth and acrossboth sides of the weld.

A 0?360?363 mm gauge volume was defined using0?3 mm wide slits on the incident and scattered sides,with a height limiter on the incident side for thelongitudinal direction. The transverse and normaldirections were measured using a gauge volume of0?360?3620 mm, with the long dimension of thisvolume parallel to the weld direction. The use of anelongated gauge volume in the welding direction wasdeemed appropriate as the residual stress is not expectedto vary significantly along the central portion of thewelded plates. The gauge volume selected was deemed tobe the smallest possible that was capable of producingdefined diffraction peaks. The centroid of the gaugevolume could therefore be positioned at 0?15 mm belowthe surface, while simultaneously avoiding partialimmersion errors. It was assumed that the principalstresses were parallel to the plate edges.

A monochromatic beam with a wavelength ofy1?66 A was provided from a squeezed Ge (004)monochromator set to a take off angle 2hM571?88u.The diffracted intensities were recorded on a positionsensitive detector. Measurements of the {211} ferritediffraction peak were performed around a scatteringangle of 2hS

The elastic strain ehkl at each measurement locationand direction was then found as follows

ehkl~dhkl{d0,hkl

d0,hkl(1)

where dhkl is the interplanar spacing in the weld for areflection (hkl).The stress sii in each of the three orthogonal

directions was found from the strain measurementsaccording to:

sii~E

1zneiiz

n

1{2ne11ze22ze33

h i(2)

where the Youngs modulus E5220 GPa and Poissonsratio n50?28, are the diffraction elastic constants for the{211}.30 i51, 2, 3 denotes the direction of the latticespacing measurement and hence, strain and stressdirection with respect to the welded plate geometry.

X-ray diffractionThe residual stresses measured using X-rays wereperformed on an Xstress 3000 instrument using Cr Karadiation on the {211} for the ferrite phase. A 2 mmcollimator was used and penetration depths were,10 mm from the sample surface, along the planeidentified in Fig. 2. Stresses in the longitudinal andtransverse orientations of the welded plate were inferredthrough the sin2 y technique.31 The y angle was variedbetween 245 and z45u (eight angles in total).Measurements were performed on the as receivedwelded plates, with no prior grinding or polishing.

Results and discussionFor each of the welded plates, three sets of stress profilesare presented. These include surface stresses measuredby neutron diffraction and X-rays. Further measure-ments were made at 2?5 mm below the surface, to allowcomparison with data collected from greater depthsbelow the free surface at other neutron sources usinglarger gauge volumes. It should be noted that the surfacestress results measured by neutron diffraction in this

work were volume averaged over 00?3 mm below thesurface. Macrographs of the three welds, as shown inFig. 4, reveal the weld layer structure. The boundarybetween passes is particularly pronounced when LTT-2is deposited on the HTT filler.

Surface stresses: neutronsThe surface residual stresses measured in three orienta-tions by neutron diffraction are presented in Fig. 5.Larger experimental uncertainties are apparent in theweld metal than the baseplate for all three specimens,which may be attributed to solidification texture.However, trends are evident and symmetry along thecentreline would suggest that the data are reliable.The surface stress profile for Weld L1 (Fig. 5a) shows

the characteristic high tensile longitudinal stresses in theHAZ with compression in the weld metal, which ischaracteristic of these types of LTT alloys and arises as aresult of the shape deformation and net expansionfollowing transformation.13 Conversely, the transversestresses are tensile in nature in the weld metal and mildlycompressive in the HAZ. The normal stresses arecompressive in the weld metal. Given their proximityto a free surface they may be expected to be closer tozero, but it is possible for stress to be retained within afew hundred micrometres of material.The filler for weld H1 (Fig. 5b) has a sufficiently high

MS that the benefits of transformation plasticity areeradicated on further cooling to ambient temperatureand this is reflected by the residual tensile longitudinalstresses in the weld metal. As with weld L1, high tensilelongitudinal stresses exist within the HAZ. The trans-verse stress is tensile in the weld metal and does notdisplay the sharp change to a compressive stress at thefusion boundary, which is evident for weld L1. Thenormal stress fluctuates about the zero stress mark inboth the weld metal and baseplate.Weld L2 (Fig. 5c), which has a capping pass of a

highly alloyed LTT filler only, displays a stress profilemore akin to weld L1. The longitudinal stress iscompressive in the weld metal with high tensile stressesin the HAZ. Continuing the trend of the previous two

4 Macrostructures of a weld L1, b weld H1 and c weld L2

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specimens, tensile transverse stresses are found in theweld metal. The normal stress is generally mildly tensile,but also fluctuates about zero.

The transverse stress measured at the critical locationof the weld toe differs significantly between weld L1(about 2600 MPa) and weld H1 (about 2200 MPa).Also, the compressive region for weld L1 extends furtherinto the baseplate and weld metal. The filler MS clearlyinfluences the transverse residual stress distribution, buttension remains in the central portion of the weld metal.This is not the case for the stresses measured in thelongitudinal orientation. A possible explanation for thisbehaviour is that the longitudinally generated stressesare the greatest because this is the direction of maximumthermal constraint during cooling. External stresses areknown to initiate variant selection during the earlystages of the cRa9 transformation.32 It is, therefore,presumed that the martensitic variants initially orientthemselves in order to cancel the dominant longitudinalstresses. However, as the transformation continues,subsequent variants are less free to align themselves insuch a manner as to minimise the tensile stresses in thetransverse orientation. This hypothesis may be con-firmed by detailed microstructural analysis; however,this is beyond the scope of this paper.

The fusion boundary at the sample surface is the siteof greatest interest due to this being the predominantlocation of failure during fatigue experiments.33 Theresidual stresses at the fusion boundary for weld L1show compression in all orientations. In contrast, thelongitudinal stresses in weld H1 are tensile in nature.The maximum tensile stresses at the fusion boundary for

weld L2 are less tensile compared with weld H1, whichwould suggest that there are benefits to applying an LTTcapping pass in multipass welds, but not to the extent ofa full LTT weld (weld L1).

Surface stresses: X-raysLongitudinal and transverse residual stresses measuredwith X-rays using the sin2 y method are presented withthe surface measurements made with neutrons overlaidin Fig. 6. The error bars from the neutron data have notbeen included for clarity, but can be referred to in Fig. 5.Some of the X-ray results had excessive errors and havebeen omitted, which may be attributed to solidificationtexture. While it is useful to compare both the X-ray andneutron data and anticipate similarities, it should benoted that measurements were made with differentsampling volumes and at slightly different depths belowthe surface.It is broadly apparent that the stresses measured by X-

rays and neutrons are in agreement. The X-ray results alldisplay tensile transverse stresses in the weld metal andthe effects of using an LTT filler show a reversal oftensile stress (Fig. 6b) to compression (Fig. 6a) in theweld in the longitudinal orientation. The X-ray mea-surements do not appear to fully replicate the peaklongitudinal tensile stresses in the HAZ, which wereidentified by neutron diffraction. Perhaps the profilesgenerated by the two techniques may show greatercorrelation if the number of X-ray measurementlocations were to be increased. However, completeagreement between the X-ray and neutron diffractionresults cannot be expected due to the difference inexperimental conditions.

a weld L1; b weld H1; c weld L25 Near surface residual stresses measured at depth of 0?15 mm by neutron diffraction

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Bulk stresses: neutronsThe small gauge volume necessary to measure surfacestresses deceases the volume of material sampled andtherefore, the number of grains. In order to verify theeffects of the reduced sampling volume, measurementswere made at 2?5 mm below the surface and comparedwith previously collected data (Fig. 7).23,34 The reduc-tion in sampling volume is significant, 1006 for weldsL1/L2 and 306 for weld H1. The trends identified forboth sets of data for the three welded plates arecomparable, the only exception being the extent ofcompressive stresses measured in the weld metal for weldL2. This is not necessarily a discrepancy because thelarger gauge volume will average the strains measured ata depth of 2?51?5 mm, while the small gauge volumerange is only 2?50?15 mm. This could have asignificant effect, depending on the stress gradients inthis region. Error bars have been included for thelongitudinal orientation and are representative of allorientations for the small gauge volume. Error bars areencompassed within the marker for the large gaugevolume measurements. Critically, the stress distributionsmeasured in the vicinity of the surface by both neutronsand X-rays are significantly different to those measuredat 2?5 mm below the surface, with high tensile transversestresses at the surface in the weld metal being the majordifferential.

ConclusionsMeasurement of the surface residual stresses producedfollowing the welding of a series of ferritic steel plateswith low transformation temperature filler alloys has

been performed using X-ray and neutron diffraction.From this study, the following conclusions and recom-mendations have been drawn:

1. The adoption of a small gauge volume haspermitted the measurement of near surface residualstresses by neutron diffraction. For the first time, thistechnique has been employed to measure the residualstresses in LTT welds and reveal the local stressvariations at the critical fusion boundary location.

2. The surface stress distributions measured byneutron diffraction are comparable with those obtainedfrom laboratory X-rays using the sin2 y technique.3. Two LTT welding alloys have been shown capable

of inducing compressive longitudinal residual stressesinto the surface layers of the weld metal for multipasswelds. Both neutron and X-ray diffraction confirm thesefindings.

4. Tensile transverse stresses were measured in theweld metal for all three welded plates. The stressesbecame compressive at the fusion boundaries, with weldL1 displaying the greatest levels of compression.

5. Weld L1 appears to display the most desirableresidual stresses as they are compressive in nature at thefusion boundary, which is the expected site of crackinitiation and subsequent propagation during service. Incontrast, weld H1, which was produced with a conven-tional filler alloy, produces tensile longitudinal stressesat the fusion boundary.

6. Weld L2, which has a singular capping pass madeby an LTT filler, is capable of inducing compressivelongitudinal stress in the weld metal near the surface andthe stress profile across the fusion boundary appearspreferable to weld H1.

a weld L1; b weld H1; c weld L26 Surface residual stresses measured using X-rays (solid lines) and neutrons (dashed lines)

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Acknowledgements

We are grateful to ESAB AB for sponsoring thisresearch. This work is based upon experiments per-formed on the L3 beam line at the Canadian NeutronBeam Centre, Chalk River Laboratories, Ontario,Canada.

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