Friction Stir Welding in HSLA-65 Steel: Part I. Influence of WeldSpeed and Tool Material on Microstructural Development

S.J. BARNES, A.R. BHATTI, A. STEUWER, R. JOHNSON, J. ALTENKIRCH,and P.J. WITHERS

A systematic set of single-pass full penetration friction stir bead-on-plate and butt-welds inHSLA-65 steel were produced using a range of different traverse speeds (50 to 500 mm/min) andtwo tool materials (W-Re and PCBN). Microstructural analysis of the welds was carried outusing optical microscopy, and hardness variations were also mapped across the weld-plate crosssections. The maximum and minimum hardnesses were found to be dependent upon bothwelding traverse speed and tool material. A maximum hardness of 323 Hv(10) was observed inthe mixed martensite/bainite/ferrite microstructure of the weld nugget for a welding traversespeed of 200 mm/min using a PCBN tool. A minimum hardness of 179 Hv(10) was found in theouter heat-affected zone (OHAZ) for welding traverse speed of 50 mm/min using a PCBN tool.The distance from the weld centerline to the OHAZ increased with decreasing weld speed due tothe greater heat input into the weld. Likewise for similar energy inputs, the size of the trans-formed zone and the OHAZ increased on moving from a W-Re tool to a PCBN tool probablydue to the poorer thermal conductivity of the PCBN tool. The associated residual stresses arereported in Part II of this series of articles.

DOI: 10.1007/s11661-012-1110-z� The Minerals, Metals & Materials Society and ASM International 2012

I. INTRODUCTION

THERE has been considerable interest in the frictionstir welding (FSW) process since it was developed byTWI in 1991.[1] FSW is a solid-state joining processderived from conventional friction welding, in which anonconsumable rotating cylindrical tool is plunged intothe material at the interface between the plates to bejoined and is then moved along the interface. Frictionalheat generated primarily from the tool shoulder, and tosome extent also from the pin, softens the materials to bejoined, which are then plastically deformed around theFSW tool in a constrained extrusion process to combineat the rear of the tool.[2] A high-quality solid-state weldedjoint can be produced, since any surface oxides arebroken up and dispersed through local plastic deforma-tion. When the process parameters are controlled cor-rectly, no parent metal melting occurs.[3] The FSW

process has been researched and developed extensivelyfor joining aluminum alloys.[4] Many aluminum alloys,particularly the high-strength ones used for aerospaceapplications, are difficult to join by conventional fusionwelding techniques. FSW can be used to join aluminumalloys without problems such as porosity or solidificationcracking and with lower distortion when compared tostandard fusion welding processes. Additionally, FSWoffers the possibility of joining dissimilar materials.[5]

Friction stir welds often exhibit improved mechanicalproperties in comparison to fusion welds of the samematerials[6] and have been successfully produced formost classes of commercial aluminum alloys.[6–12]

Although joining of aluminum-based alloys has beenthe main focus of FSW development, the process can beused for other nonferrous and ferrous alloys. Untilrecently, relatively little research had been published onthe FSW of steels,[13–30] but the literature on the subjectis expanding rapidly. There have been two mainobstacles to the development of FSW for joining steels.First is the development of inexpensive tool materialscapable of surviving the high temperatures and forcesgenerated by the process.[31–34] Second, the fact thatfusion welding processes are already established forjoining steels (and are relatively inexpensive, not requir-ing large machines and complicated setup) reduces thedriving force for development. It is clear, however, thatif FSW could be developed for ferrous alloys, then itoffers the possibility of joining thinner section, higherstrength alloys using automated procedures giving lowerdistortion than conventional arc welding processes.[17]

This could be of great interest in, for example, the ship-building industry, in which high-strength low-alloy(HSLA) steels are used extensively.

S.J. BARNES, NDT Manager, and P.J. WITHERS, Professor ofMaterials Science, are with the School of Materials, Materials ScienceCentre, University of Manchester, Manchester M13 9PL, UnitedKingdom. Contact e-mail: simon.barnes-2@manchester.ac.uk A.R.BHATTI, formerly Research Assistant, with the School of Materials,Materials Science Centre, University of Manchester, is now PostdoctoralResearch Assistant, with the Department of Materials Science andMetallurgy, University of Cambridge, Cambridge CB2 3QZ, UnitedKingdom.A. STEUWER,Visiting Professor forMaterials Engineering, iswith ESS Scandinavia, University of Lund, 22350 Lund, Sweden, and isalso with NMMU, Port Elizabeth 6031, South Africa. R. JOHNSON,formerly Project Leader, with TWI Yorkshire, Wallis Way, Catcliffe,RotherhamS605TZ,UnitedKingdom, isnowretired. J.ALTENKIRCH,formerly Researcher, with the School of Materials, Materials ScienceCentre, University of Manchester, is now Project Manager, with SiemensAG, Mlheim an der Ruhr D-45473, Germany.

Manuscript submitted July 14, 2010.Article published online April 11, 2012

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Some advances have been made in the development ofFSW for steels since information about the techniquewas first published in 1998.[35] These advances have beenmainly through improved materials selection and tooldesign.[31–34] Tool material development work haslooked primarily at refractory metals such as tungsten-rhenium (W-Re),[13,18,33] polycrystalline boron nitride(PCBN),[33,36] and tungsten carbide (WC),[26,27,34] butnone of these solutions is yet fully optimized or cost-effective. The use of PCBN tools has shown promisesince the material has relatively good wear and temper-ature resistance, but it has limited ductility and can beprone to pin failure if high traverse loads are placedupon it.[36] This effectively limits the section thicknessesthat can be joined and the welding traverse speeds thatcan be used. Refractory metals, particularly W-Re, onthe other hand, offer significantly greater ductility thanPCBN and so can generally be used at higher weldingtraverse speeds, but these tools currently can suffersignificant tool wear, which limits the effective tool lifeand therefore increases cost. Work is underway toimprove the wear resistance of W-Re through theaddition of cubic boron nitride and diamond.[33]

Although many tool materials have been tested, therehas been relatively little work published[18,21,22,26,27] onthe effect of systematic variations in process parameterson the microstructure and properties of the friction stirwelded joints produced with these tools. This work isrequired to enable a greater understanding of theprocess window for the production of defect-free weldswith acceptable microstructures in high strength lowalloy steels. It has been shown[37] that the cooling ratesobserved in steel FSWs are comparable to arc weldingprocesses with heat inputs of 1.3 kJ/mm. Thus, it mightbe expected that the transformations produced by FSWshould be similar to conventional fusion welding pro-cesses. However, it should also be borne in mind that thepeak temperatures reached are considerably lower,which might be expected to influence austenite grainsize and thus hardenability.[37] In addition, the highlevels of deformation associated with the FSW processcould influence recrystallization in the austenite phasefield. Given that the peak temperatures attained, levelsof deformation, and local cooling rates are all expectedto be influenced greatly by the welding traverse speed,rotation speed, and tool material, it is necessary tounderstand better the relationship between FSW condi-tions and weld microstructure, properties, and residualstresses.

Consequently, the objective of the present study wasto examine in a systematic manner the effect of weldingtraverse speed and tool material on the microstructureand hardness distribution across friction stir welds in thecommercially important HSLA steel HSLA-65. In viewof the critical importance of residual stress for welddistortion and in-service performance, the dependence

of the residual stress distribution across the samesystematic set welds was also examined. These resultsare reported in a companion article.[38]

II. EXPERIMENTAL PROCEDURE

A. Materials and Welding Details

HSLA-65 is a HSLA steel with nominal yield strengthof 450 MPa and a relatively high level of toughness dueto the low carbon content. The chemical composition ofthe HSLA-65 steel plate used in the welding trials isshown in Table I. 600-mm 9 110- to 120-mm plates(butt welds (BWs)) and 600-mm 9 140-mm plates(bead-on-plate (BOP) welds) of 6.35-mm (1/4-in.) thick-ness were processed at TWI (Rotherham, United King-dom) by friction stir butt welding using a PCBN tooland BOP welding using a W-Re tool (nominal compo-sition W-25 wt pct Re). In the former case, the toolgeometry was a 30 deg cone angle tapered pin designwith a 20 TPI stepped spiral cut into the surface. Thepin was 5.5 mm in length pin with a 23.7-mm-diameterspiral convex shoulder. The W-Re tool had similardimensions to the PCBN tool, but had a TRI-FLUTE*

pin design with a pin diameter of 8 mm at the shoulderand 6 mm at the tip mounted on a 25-mm-diameterconcave shoulder. Photographs of the tool geometriesare shown in Figure 1. The welding parametersemployed for these samples are summarized in Table II.The range of traverse speeds examined for the W-Retool were from 50 to 500 mm/min, whereas the PCBNtool was only used up to a traverse speed of 250 mm/min because higher speeds were seen to cause pinfracture in pretrial welds. PCBN and W-Re were chosensince they showed the most promise as tool materials forFSW of steel, but each tool material had weaknesses, asalready discussed in Section I. The tool designs werechosen since these showed the most consistent weldquality with each tool material when welding HSLA-65steel.All of the welds were carried out parallel to the rolling

direction of the steel plate. The heat input per unitlength, E, of each weld (in units of kJ/mm, sometimescalled ‘‘line energy’’) can be estimated using Eq. [1]:

E ¼ x � s=1� 106 � m ½1�

where x is the angular velocity of the spindle (radiansper second), s is the applied torque (Nm), and v is thetraverse speed of the weld tool (m/s). It should be notedthat Eq. [1] only gives an estimate of the heat input into

Table I. Nominal Composition (Weight Percent) of the HSLA-65 Steel Plate (Fe Balance)

C Si Mn P Cr Mo Nb Ti V Ni Cu Al S

0.09 0.26 1.42 0.06 0.15 0.07 0.02 0.01 0.06 0.35 0.26 0.02 0.01

*TRI-FLUTE is a trademark of TWI Ltd., Granta Park, GreatAbington, Cambridge, UK.

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the weld, since the equation takes no account of heatlosses through radiation or by conduction from the tooland shoulder to the tool holder, which will bothinfluence the weld temperature attained for a specificheat input. Ideally, the temperature in the weld zoneshould be measured as a function of heat input to

understand the relative contributions of the differentheat loss mechanisms. In the current work, no specificweld temperature measurements were taken. The possi-ble changes to weld temperature were only inferred fromthe microstructural changes observed in the welds as afunction of weld process changes.Figure 2 shows that, although the rotation speeds and

applied forces summarized in Table II for the two toolmaterials were different, the line energies of the two setsof welds overlap.

B. Microstructure Analysis and Hardness Measurements

All the welds were sectioned, metallographicallyprepared, and etched in 2 pct Nital. Optical photomi-crographs of selected regions of the weld microstruc-tures were obtained using a Zeiss Axioimager M1moptical microscope. The distances from the weld center-line to the edge of the weld nugget and the outer heat-affected zone (OHAZ) were measured accurately using aZeiss Axiovision image analysis system at various depthsbelow the top surface of the welded plates. The weldmicrosections were also examined in a Zeiss scanning

Fig. 1—Photographs showing the PCBN and W-Re tool designsused in the welding trials.

Table II. Welding Parameters for the HSLA-65 BOP and BWs Studied

ToolMaterial

RotationSpeed (rpm)

TraverseSpeed (mm/min)

End (X)Force (kN)

Down (Z)Force (kN)

Torque(Nm)

Weld Energy(kJ/mm)

WeldType

W-Re 600 50 0.3 12 49 3.69 BOPW-Re 600 100 0.4 15 58 2.19 BOPW-Re 600 150 0.4 18 65 1.63 BOPW-Re 600 200 0.7 21 68 1.28 BOPW-Re 600 250 3.6 23 76 1.15 BOPW-Re 600 300 6.5 31 87 1.09 BOPW-Re 600 350 9 45 98 1.06 BOPW-Re 600 400 10 45 102 0.96 BOPW-Re 600 450 11 51 106 0.89 BOPW-Re 600 500 15 55 113 0.85 BOPPCBN 400 50 6 35 62 3.12 BWPCBN 400 100 5 40 78 1.96 BWPCBN 400 150 5 42.5 91 1.52 BWPCBN 400 200 9 47.5 98 1.23 BWPCBN 400 250 12 52.5 111 1.12 BW

Fig. 2—Graph showing the line energy (welding heat input per unitlength) calculated for the W-Re tool BOP welds and the PCBN toolBWs as a function of welding traverse speed.

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electron microscope (Carl Zeiss MicroImaging GmbH,Gottingen, Germany) using both secondary and back-scattered imaging modes. An Oxford Instruments (HighWycombe, Buckinghamshire, UK) energy dispersiveX-ray analysis system was used to determine the chemicalcomposition in different areas of the weld.

Vickers hardness traverses were carried out on someof the W-Re tool welds at the Corus Research Devel-opment and Technology Centre (Swinden TechnologyCentre, Rotherham, United Kingdom), using a 10-kgload. Hardness values were obtained at 1-mm steps to adistance of 16 mm on either side of the weld centerline.These measurements were performed at depths of 1.3,2.3, 3.3, and 4.3 mm from the top surface, to enablecontour maps of the hardness levels to be drawn.Additionally, Vickers hardness measurements were per-formed on the PCBN tool welds at FaME 38, ILL-ESRF (Grenoble) also using a 0.5-kg load with apyramidal diamond indenter on a Buehler OmniMetsystem. These measurements were also obtained at 1-mm steps to a distance of 16 mm each side of the weldcenterline. These measurements were performed atdepths of 1, 2, 3, 4, and 5 mm from the top surface to

enable contour maps of the hardness levels to be drawnfor these welds.

III. EXPERIMENTAL RESULTS

A. Influence of Welding Parameters on Microstructure

Macro- and micrographs of the weld cross sections inHSLA-65 steel for welding traverse speeds of 50 and500 mm/min using a W-Re tool are shown in Figures 3and 4, respectively. The macrostructure of the weld canbe split into several distinct regions in general agreementwith published work.[39] These are the weld nugget (N);a region either side of the weld nugget termed thethermomechanically affected zone (TMAZ); the inner,middle, and outer HAZ regions (IHAZ, MHAZ, andOHAZ, respectively); and the parent metal (P). Repre-sentative micrographs were obtained from each of theseareas of the weld.Figure 5 shows a comparison of the microstructures

in the nugget center and OHAZ regions for fourdifferent welding traverse speeds using the W-Re toolmaterial. It is clear from these micrographs that the

Fig. 3—Typical macrostructures and microstructures for a single-pass partial penetration friction stir weld in HSLA-65 steel made using a W-Retool at a welding speed of 50 mm/min. Areas marked A and R represent advancing and retreating sides of the weld, respectively. The areas ofthe weld shown are parent metal (P), weld nugget at 1.5-mm depth (N(1.5)), weld nugget at 3-mm depth (N(3)), IHAZ, MHAZ, and OHAZ.Scale bar represents 20 lm in each image.

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traverse speed has a significant influence on microstruc-ture in both the weld nugget and OHAZ.

Figure 6 shows the influence of welding speed, weldingheat input (line energy), and tool material on the extent ofthe transformed zone (i.e., stir zone plus TMAZ plusIHAZ) and theOHAZatmidthickness of theweldedplate.The extent of the zones was measured using an opticalmicroscope to distinguish the edge of each zone. There isclearly some uncertainty in the measurement of the zoneedge positions, since the exact position requires expertinterpretation of the microstructure. In an attempt tounderstand the variability, the zone positions were mea-sured independently by three different operators and theaverage position calculated. The variability in the mea-surement of the zone edge positions was less than 0.5 mmfor each weld. In line with expectation, both the trans-formed zone and the OHAZ increase in size as the weldingtraverse speed decreases since there is a greater amount ofheat put into the material at lower welding speeds.

B. Influence of Welding Parameters on Weld Hardness

Maps of the measured hardness across the weldedsections for the different traverse speeds using the W-Re

and PCBN tools are shown in Figures 7 and 8,respectively. The mean hardness values measured inthe weld nugget are plotted as a function of weldingtraverse speed and weld energy for the different toolmaterials in Figures 9 and 10. In all cases, the weldnugget is harder than the parent metal and the minimumhardness area is the outer HAZ of the weld. Theposition of the minimum hardness value moves closer tothe weld centerline as the traverse speed is increased.

C. Influence of W-Re Tool Wear on Microstructure

Figure 11 illustrates the level of tool wear typicallyobserved for the FSW process on a HSLA-65 substrateusing a W-25 pct Re tool material. It can be seen thatafter 5 m of welding, all the features on the tool andshoulder have been worn completely away. This abra-sive wear of the tool results in areas rich in tungstenbeing deposited in the weld. Figure 12 shows an exampleof this deposition in the 50-mm/min weld. Quantitativeenergy dispersive X-ray analysis results shown inFigure 12 indicate that bands of material richer in tungstenare present at the edge of the weld nugget. These bandswere difficult to image clearly using optical microscopy,

Fig. 4—Typical macrostructure and microstructures for a single-pass partial penetration friction stir weld in HSLA-65 steel made using a W-Retool at a welding speed of 500 mm/min. Areas marked A and R represent advancing and retreating sides of the weld, respectively. The areas ofthe weld shown are parent metal (P), weld nugget at 1.5-mm depth (N(1.5))), weld nugget at 3-mm depth (N(3)), TMAZ, IHAZ, and OHAZ.Scale bar represents 20 lm in each image.

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but SEM examination using both secondary and back-scattered imaging modes was more successful. Thetungsten-rich bands appear darker than the main area

of the weld nugget using secondary electron imagingand, as expected, brighter using backscattered electronimaging due to the high atomic number of tungsten.

Fig. 5—Photomicrographs showing (a) through (d) the weld nugget microstructure and (e) through (h) the OHAZ microstructure at 1.5-mmdepth for welds produced using a W-Re tool. Scale bar represents 20 lm in each image.

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Figure 13 shows higher magnification views of the darkband region. In this area, large numbers of fine scale (ofthe order of 1 lm) particles were observed. These finescale particles were found to contain W and Re.Although all the quantitative X-ray results shown inFigures 12 and 13 were obtained after ZAF correction,it should be noted that the excitation volume for X-raysbelow the surface of the sample (in this case approxi-mately 1 lm depth for an accelerating voltage of20 keV) gives some uncertainty in the compositionmeasurement for the phases seen on the sample surface,since their thickness is unknown. It can only beconcluded that the dark bands observed at the edge ofthe weld nugget optically in the 50-mm/min weldcontained greater amounts of W and Re than otherareas of the weld.

IV. DISCUSSION

A. Weld Microstructure

The original parent metal microstructure (microstruc-ture P in Figures 3 and 4) of the HSLA-65 steel is typicalof low alloy steel having a structure consisting of ferriteand pearlite, with a fine network of niobium carbidesserving to pin the grain boundaries and thereby limitgrain growth during processing of the plate. As notedpreviously,[30] the HSLA-65 parent metal exhibits a fineequiaxed ferrite structure with some banding of the

pearlite structure parallel to the rolling direction of theplate. The local postweld microstructure is a function ofthe thermal and mechanical cycles experienced duringthe welding process, as shown schematically in Figure 14.While no melting occurs during FSW, very hightemperatures were measured,[14,25] certainly in excessof the transformation temperature, A3. On the basis ofprevious work,[40] the transformation temperatures forHSLA-65 were calculated as A1 = 1000 K (727 �C) andA3 = 1136 K (863 �C). Although no specific weldtemperature measurements were performed in the cur-rent study, it is highly likely, on the basis of previouswork[14,21,22,25,27] and examination of the weld micro-structures, that, in the center of the weld nugget regionfor all traverse speeds, the metal is heated into theaustenite phase field (above A3) and is plasticallydeformed around the FSW tool before experiencingrapid cooling once the tool has passed. However, it mustbe noted that, in the current study, the examination ofthe weld microstructure and hardness only allowsrelative changes in peak weld temperatures and coolingrate, as a function of welding process parameters, to beinferred. It is not possible to state that a certaintemperature was obtained with certainty in the absenceof specific weld temperature measurements. Even whenweld temperature measurements are obtained, someextrapolation is usually required, since either the surfacetemperature is measured using infrared thermographyor thermocouples are embedded in the plates to bewelded slightly away from the tool traverse path toavoid thermocouple damage.In the case of the HSLA-65 steel, microstructural

examination indicates that this thermal excursion leadsto a mixed martensite, bainite, and proeutectoid ferritemicrostructure in the center of the weld nugget (N inFigures 3 and 4). This result is in agreement withprevious work on the FSW of DH36 steel,[18] RQT-701steel,[24] and the CCT curve for HSLA-65.[40] InFigures 3 and 4, it can be seen that the prior austenitegrain size and the acicular nature of the ferrite in the nuggetmicrostructure decrease with increasing depth below theplate surface. This is expected since the heating effect ofthe tool shoulder decreases with increasing depth intothe plate. Thus, the peak temperature is expected todecrease with increasing depth, resulting in a reductionin austenite grain size and an increase in the temperaturerange for ferrite growth during cooling. The mixedmartensite, bainite, and ferrite microstructure extendsright across both the weld nugget and the stirredmicrostructure (TMAZ) region immediately adjacentto the FSW tool. As noted in previous work,[30] thepresence of the TMAZ in HSLA-65 welds can be quitedifficult to determine without special etching techniques,but in this case, a thin region toward the edges of thenugget showing a finer transformed microstructure(Figure 4) was thought to be the TMAZ, but furtherinvestigations are required to confirm this.The phase fractions in the weld microstructure are

difficult to quantify and depend upon the peak temper-ature (Tp) reached during welding and the cooling rateafter welding (characterized by Dt8/5—the cooling timebetween 1073 K (800 �C) and 773 K (500 �C)[40]), but

Fig. 6—Graphs showing the distance from the weld centerline to theedge of the transformed zone (nugget + IHAZ) and the outer HAZas a function of welding speed and line energy for different toolmaterials. The distances were measured at the center depth of thewelds on the advancing side.

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the presence of this combination of phases is expected.Research on the CCT diagram for HSLA-65 steel[40]

found mixtures of martensite, bainite, and ferrite forcooling rates having Dt8/5 values ranging from 1.5 to19 seconds. Previous work[37] on FSW of steel identifieda cooling rate (Dt8/5 ~11seconds) similar to a traditionalarc welding process at 1.3 kJ/mm energy.

The IHAZ, MHAZ, and OHAZ regions exhibit verydifferent microstructures consistent with attaining dif-ferent maximum temperatures during welding(Figure 14). The IHAZ exhibits a mixed martensite,bainite, and ferrite microstructure similar to the centerof the weld nugget, but with a finer prior austenite grainsize and more polygonal ferrite consistent with reachinga lower maximum temperature (but still above A3) thanthe center of the weld nugget. The MHAZ delineates anintercritically heated region, where some austeniteformation occurs during the FSW process and then

transforms to bainite and ferrite on cooling. The outerHAZ consists of an overtempered martensite structurewith ferrite and carbide phases present. This area wasnot transformed during welding, but it was heated to atemperature (>923 K (650 �C)) sufficient to cause sig-nificant coarsening and spheroidization of the carbidesto occur.It is evident in Figure 5 that the traverse speed has a

significant influence on microstructure in both regions.As previously discussed, when considering the effect ofwelding traverse speed on microstructure, it is necessaryto consider the peak temperature attained during FSWand the cooling rate after the welding tool has passed.For example, at the slowest welding speed (50 mm/min),a high peak temperature would be expected, but aslower cooling rate after welding would also occur dueto the lower temperature gradient between the center ofthe weld and the rest of the steel plate. In the case of the

Fig. 7—Hardness contour maps in Hv (10) for single-pass friction stir welds in 6.35-mm-thick HSLA-65 steel plates made using a W-Re tool atvarious traverse speeds with a 600 rpm tool rotation speed. Advancing side of the weld is on the right side.

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fastest welding speed (500 mm/min), the reverse wouldbe expected to be the case—a lower peak temperaturebut a faster cooling rate after welding. In the case of theweld nugget, the material is heated into the austenitephase field (above A3) for all the weld speeds, but boththe degree of superheating and the time at temperaturewill depend upon the welding traverse speed. The higherpeak temperatures expected from lower welding speedsgive larger austenite grains. This grain size then influ-ences the hardenability of the material and thus thephase type and morphology that forms on cooling.These issues were discussed previously for fusion weldsin low alloy steels.[41] Increases in the original austenitegrain size lead to a promotion of the formation ofacicular ferrite due to a lowering of the start tempera-ture for polygonal ferrite formation (Pfs) accompanied

by a small increase in the start temperature for acicularferrite formation (Afs). Additionally, with increases inaustenite grain size, more displacive transformationproducts (i.e., martensite and bainite) can be formed at aparticular cooling rate. Increases in cooling rate, on theother hand, will lead to a higher likelihood of displacivetransformation products in the weld nugget.The changes in weld nugget microstructure with

increasing traverse speed shown in Figure 5 can beexplained in terms of these issues. On moving from 50 to100 mm/min (Figure 5(a) and (b)), there is a significantincrease in the amount of martensite, bainite, andacicular ferrite in the nugget microstructure due to theincrease in cooling rate after welding. On moving from100 to 300 mm/min (Figure 5(b) and (c)), the influenceof the peak temperature starts to override the cooling

Fig. 8—Hardness contour maps in Hv (10) for single-pass friction stir welds in 6.35-mm-thick HSLA-65 steel plates made using a PCBN tool atvarious traverse speeds with a 400 rpm tool rotation speed. Advancing side of the weld is on the right side.

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rate effect. Although the cooling rate after weldingshould increase with increasing welding speed, the 300-mm/min weld microstructure (Figure 5(c)) shows morepolygonal ferrite and less displacive transformationproducts than observed for the 100-mm/min weld(Figure 5(b)). This is due to the lower peak temperatureand thus smaller austenite grain size observed in the300 mm/min sample. This increases the temperaturerange for polygonal ferrite growth and gives smallermicrostructure regions for the displacive transformationproducts to form on cooling. On moving from 300 to500 mm/min (Figures 5(c) and (d)), the cooling rateeffect starts to become more dominant again with moreacicular ferrite being formed on cooling. The OHAZmicrostructure is also influenced by changes to thewelding traverse speed, since this influences the amountof heat put into the material and the cooling rate in theOHAZ area. It can be seen in Figure 5 that the OHAZmicrostructure is coarsest for the 50 mm/min sample(Figure 5(e)) in line with expectations, since at thiswelding speed there is the greatest heat input into theweld and the slowest cooling rate in the OHAZ area.This leads to the greatest degree of overtempering of theparent metal microstructure in this weld.

Figure 6 shows that there is an essentially lineardependence of the transformed zone size on weldingspeed, although of shallow gradient, indicating a rela-tively weak effect of peak temperature with increasingweld speed. The size of the OHAZ is also essentiallylinear with welding speed except at low welding speeds,

where the size of the OHAZ increases markedly. This isdue to the increase in energy transferred to the weld atlower traverse speeds combined with lower cooling ratesafter welding, giving rise to greater volumes of materialheated to sufficiently high temperatures and times toallow significant coarsening of the microstructure. Theeffect of line energy is also shown in Figure 6. As the lineenergy increases, the sizes of both the transformed zoneand the OHAZ increase, as expected, due to the greateramount of heat transferred to the weld. Figure 6 alsoallows a comparison of tool materials. In this respect,there appears to be little influence of tool material whenlooking at the relationship between zone sizes andtraverse speeds for both W-Re and PCBN tools, but itshould be remembered that different rotation speedswere used for these tool materials (Table II). Conse-quently, the relationship between zone size and lineenergy represents a fairer comparison. Although thereare only five results for the PCBN tool, it can be seenthat the transformed zone and OHAZ sizes are slightlysmaller for the PCBN tool. This is an expected result onthe basis of the line energy comparison when thedifferent thermal properties of the two tool materialsare considered. PCBN is a better thermal conductorthan W-Re and is thus more effective in conducting heataway from the weld zone during welding. Sampleswelded using a PCBN tool therefore would be expectedto attain lower temperatures than for a W-Re tool forthe same energy input. Thus, it would be expected thatthe PCBN welds would exhibit smaller transformed

Fig. 9—Graphs showing the mean hardness and variability in theweld nugget at center depth or the W-Re tool as a function of (a)welding traverse speed and (b) line energy.

Fig. 10—Graphs showing the mean hardness and variability in theweld nugget at center depth for the PCBN tool as a function of (a)welding traverse speed and (b) line energy.

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zone and OHAZ sizes than the W-Re tool welds.Although the energy inputs are similar for similar weldtraverse speeds using W-Re and PCBN tools, it can beseen in Table II that the forces used for the PCBN weldsare significantly greater particularly at low traversespeeds. This implies that the PCBN weld is at a lowertemperature than the W-Re weld, since more force isrequired to move the tool. TWI generally use lowerrotation speeds for the PCBN tools to keep the tooltemperature lower, since this reduces tool degradationand subsequent cracking. In addition, it should beremembered that the tool designs for PCBN and W-Retools were also different, and so the difference betweenthe tools is not solely a result of the tool material. TheTRI-FLUTE** design of the pin in the W-Re tool

would be expected to give better material flow up anddown the pin, and thus generate greater frictionalheating, than the tapered thread pin design of thePCBN tool. This would exacerbate the differences inweld temperature expected from the thermal conductiv-ity differences of the two pin materials, althoughincreasing tool wear in the case of the W-Re tool willreduce the effect of the tool design differences.

B. Weld Hardness

Figures 7 and 8 show that, in all cases, the weldnugget is harder than the parent metal. This is expectedfrom the transformed nature of the nugget microstruc-ture. The hardness across the welds exhibits a minimumvalue associated with the outer HAZ of the weld. This isalso expected from the microstructural examination,since the outer HAZ is overtempered, resulting insoftening as a result of grain coarsening and carbidespheroidization. The position of the minimum hardnessvalue moves closer to the weld centerline as the traversespeed is increased presumably due to the lower amountof heat transferred to the weld.It can be seen in Figure 7 that the highest weld nugget

hardness is associated with the 100 mm/min traversespeed for the W-Re tool. This was expected from theweld nugget microstructure examination, since thissample contains the greatest amount of martensite. Asthe weld speed increased above 100 mm/min, the cool-ing rate would be expected to increase with a corre-sponding tendency for more martensite formation.Additionally, however, the peak weld temperaturewould also be expected to decrease with increasingwelding speed. As noted in previous work[41] this gives alower prior austenite grain size, which leads to a loweroverall proportion of martensite and a higher propor-tion of polygonal ferrite in the final microstructure. Asmall increase in maximum hardness on moving from

Fig. 11—Photographs showing (a) new and worn spiral convex shouldered FSW tools, (b) Tool after 4 welds (approximately 1m total weldlength), (c) Tool after 9 welds (approximately 2.25 m total weld length), (d) Tool after 23 welds (approximately 5.75 m total weld length).

**TRI-FLUTE is a trademark of TWI Ltd., Granta Park, GreatAbington, Cambridge, UK.

2352—VOLUME 43A, JULY 2012 METALLURGICAL AND MATERIALS TRANSACTIONS A

400 to 500 mm/min traverse speed is thought to be theresult of the increased cooling rate outweighing theeffect of decreased peak temperature to give a moreacicular microstructure.

A comparison of Figures 9 and 10 shows the influenceof tool material. Although not all the same traversespeeds were examined for both tool materials, it appearsthat the peak hardness position has shifted to greatertraverse speeds for the PCBN tool. It has already beennoted that the use of the PCBN tool and processparameters is likely to give a lower temperature in thenugget material than the W-Re tool material for thesame energy input. This means that lower peak temper-atures and faster cooling rates are likely in the case of

the PCBN tool. Thus, for the same energy input, a lowerprior austenite grain size would be expected for thePCBN tool weld due to the lower peak temperature.This increases the driving force for polygonal ferriteformation with less displacive transformation products.In addition, however, the lower temperature of the toolwill also give a faster cooling rate and thus greateramounts of martensite and bainite would be expected toform on cooling for the PCBN tool weld. The compe-tition between the effect of peak temperature andcooling rate drives microstructure formation in the weldnugget. Clearly, there is a different balance betweenthese effects in the PCBN and W-Re tool welds. Itappears that the effect of the faster cooling rate

Fig. 12—Optical and SEM micrographs showing tungsten-rich bands close to the edge of the weld nugget in the 50-mm/min traverse speed weld.(a) Macrostructure of the weld. (b) Secondary electron image showing dark bands at the edge of the weld nugget. (c) Backscattered electronimage showing light bands at the edge of the weld nugget. (d) Higher magnification secondary electron image of the dark bands. Quantitative EDXanalysis results show that the dark band is rich in tungsten.

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dominates to higher line energy values in the case of thePCBN welds.

C. Tool Wear

As a result of tool wear during the welding process,tungsten-rich bands were observed in all the welds,although, qualitatively, the 50 mm/min weld appearedto have the largest areas indicating that increasing thetool temperature is detrimental to the abrasive wear ofthe tool. Further work is required to ascertain if thepresence of these tungsten-rich bands is detrimental tothe mechanical properties of the weld, but clearly thislevel of tool wear is a serious issue for industrial

application of the FSW process for steels since the costof the W-Re tool is already high. Further work isalready underway to improve the wear resistance of theW-Re alloy. The abrasive wear observed for the PCBNtool was significantly less than the W-Re tool. In thiscase, cracking of the shoulder to pin intersection leadingto tool breakage is generally a bigger problem thanabrasive wear. The low fracture toughness of the PCBNtool material makes it susceptible to mechanical crack-ing during the initial plunge stage and thermal crackingas a result of differences of heating and cooling ratesbetween the pin and shoulder. This problem results in alower limit on the traverse and rotation speeds for thePCBN tools when compared to the W-Re tools.

V. CONCLUSIONS

Good defect-free welds were obtained in all cases. Theweld nugget microstructure for the HSLA-65 steel wasfound to be a mixture of martensite, bainite, and pro-eutectoid ferrite in all the welds examined. The propor-tions of the different phases, and thus the weld nuggethardness, were influenced by the weld tool traversespeed, rotation speed, and material. Changing the weldtraverse speed influenced the nugget microstructure,presumably through its effect on the peak temperatureattained during welding and the cooling rate once thetool had passed. The actual weld temperatures were notquantified during this study, but the changes in weldmicrostructure observed as a result of changes towelding parameters were used to infer the probablevariations in weld temperature.In the case of the W-Re tool, the highest mean nugget

hardness (287 Hv(10)) was observed at a traverse speed

Fig. 13—High-magnification secondary electron images of the dark band showing large numbers of fine scale (approximately 1 lm) particles.EDX analysis of these particles shows them to be rich in W and Re.

Fig. 14—Schematic showing different microstructural regions of asingle-pass friction stir weld in HSLA-65 steel. The weld thermalcycles for each area are also shown.

2354—VOLUME 43A, JULY 2012 METALLURGICAL AND MATERIALS TRANSACTIONS A

of 100 mm/min, since this gave the best combination ofprior austenite grain size and cooling rate to give thehighest amount of displacive transformation products inthe microstructure. The observed mean weld nuggethardness was found to be relatively low (248 Hv(10)) forthe lowest welding speed (50 mm/min) due to a combi-nation of high prior austenite grain size and low coolingrate. The weld nugget hardness was also low(233 Hv(10)) for a welding speed of 400 mm/min dueto a combination of lower prior austenite grain size andhigher cooling rate, which resulted in a lower overallproportion of martensite and a higher proportion ofpolygonal ferrite in the final microstructure

The minimum hardness values for each weld were179 Hv(10) for the PCBN tool material and 185 Hv(10)for W-Re tool material. In both cases, the lowesthardness was found in the lowest traverse speed weld(50 mm/min) and was associated with the OHAZ sincethis comprised over-tempered parent metal. The dis-tance from the weld centerline to the OHAZ increasedwith decreasing weld speed due to the greater heat inputinto the weld.

The tool material and design also influenced the weldmicrostructure and hardness. For similar energy input,the size of the transformed zone and the OHAZdecreased on moving from a W-Re tool to a PCBNtool. This is the result of the PCBN weld being at alower temperature than the W-Re weld for similar lineenergy values. Accordingly, the traverse speed giving thehighest hardness was also higher for the PCBN tool(200 mm/min in the PCBN weld compared to 100 mm/min in the W-Re weld), since the faster cooling rateassociated with higher traverse speeds dominated overthe peak temperature effect in the lower temperaturePCBN welds. The associated variation in residualstresses for the two sets of welds is reported in acompanion article.[38]

ACKNOWLEDGMENTS

The authors acknowledge the work of Lyn Drewett,Corus RDandT Centre. EPSRC is acknowledged forfunding, including the FSW program of the LightAlloys Portfolio Partnership (Grant No. EP/DO29201/1).

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