Microstructure and mechanical properties of friction stir welded SAF

7/24/2019 Microstructure and mechanical properties of friction stir welded SAF

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Materials Science and Engineering A 397 (2005) 376384

Microstructure and mechanical properties of friction stir welded SAF2507 super duplex stainless steel

Y.S. Sato a,, T.W. Nelson b, C.J. Sterling b, R.J. Steel c, C.-O. Pettersson d

a Department of Materials Processing, Graduate School of Engineering, Tohoku University, 6-6-02 Aramaki-aza-Aoba,

Sendai 980-8579, Japanb Mechanical Engineering Department, Brigham Young University, 435 CTB, Provo, UT 84602, USA

c MegaStir Technologies, 275 West 2230 North, Provo, UT 84604, USAd Sankvik Materials Technology, Sandviken S-811 81, Sweden

Received 18 January 2005; received in revised form 18 February 2005; accepted 25 February 2005

Abstract

Themicrostructure and mechanical propertiesof friction stir(FS) welded SAF2507 superduplexstainless steel wereexamined.High-quality,

full-penetration welds were successfully produced in the super duplex stainless steel by friction stir welding (FSW) using polycrystalline

cubic boron nitride (PCBN) tool. The base material had a microstructure consisting of the ferrite matrix with austenite islands, but FSW

refined grains of the ferrite and austenite phases in the stir zone through dynamic recrystallisation. Ferrite content was held between 50 and

60% throughout the weld. The smaller grain sizes of the ferrite and austenite phases caused increase in hardness and strength within the stir

zone. Welded transverse tensile specimen failed near the border between the stir zone and TMAZ at the retreating side as the weld had roughly

the same strengths as the base material.

2005 Elsevier B.V. All rights reserved.

Keywords: Duplex stainless steel; Friction stir welding; Mechanical properties; Microstructure

1. Introduction

Duplex stainless steels have a mixed microstructure con-

sisting of ferrite (bcc) and austenite (fcc) phases[1,2].When

duplex stainless steels have the optimum phase balance,

which is usually approximately equal proportions of ferrite

and austenite phases, they exhibit higher resistance to stress

corrosion cracking and higher strength than austenitic stain-

less steels[1].Taking advantages of these positive factors,

duplex stainless steels are widely used in the oil and gas,

petrochemical, pulp and paper, and pollution control indus-

tries[1].

It is well known that the duplex stainless steels exhibit

good weldability[1], but the melting and solidification as-

sociated with fusion welding processes destroy the favorable

duplex microstructure of these stainless steels. Microstruc-

Corresponding author. Tel.: +81 22 795 7353; fax: +81 22 795 7352.

E-mail address:[email protected] (Y.S. Sato).

ture of the wrought duplex stainless steels has a pronounced

orientation of austenite islands in the ferrite matrix, paral-

lel and transverse to the rolling direction, but fusion welding

produces a microstructure consisting of coarse ferrite grains,

and both intergranular and intragranular austenite phases in

the weld metal and heat affected zone (HAZ)[15].In gen-

eral, the volume fraction of ferrite is much higher than that of

austenite in the weld metal and HAZ. These changes in mi-

crostructure cause the loss of low-temperature notch tough-

ness and corrosion resistance in the weld[1,2]. To alleviate

these problems,careful control of theweld metal composition

and temperature are often required during welding.

Friction stir welding (FSW) is a solid-state joining process

developed and patented by The Welding Institute (TWI) in

UK in 1991[6]. Since inception, FSW had been restricted to

the lower melting temperature materials, such as aluminum

(Al) and magnesium (Mg) alloys[716]. However, over the

past 5 years, much progress has been made in FSW of high

temperature materials by numerous investigators [1725].

0921-5093/$ see front matter 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.msea.2005.02.054


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Y.S. Sato et al. / Materials Science and Engineering A 397 (2005) 376384 377

These studies have reported that FSW achieves similar grain

refinement in the stir zone of the steels as those observed in

aluminum. Additionally, FSW does not accompany melting

and solidification, alleviating the formation of porosity and

adverse phase transformations during the welding process.

This situation suggests that the favorable duplex microstruc-

ture of the duplex stainless steel should be maintained, butthere have been few papers dealing with FSW of duplex stain-

less steel.

The present study used SAF 2507 super duplex stainless

steel as the base material. This particular super duplex stain-

less steel exhibits significantly higher corrosion resistance

than early grades of duplex stainless steels[1]. The nominal

chemical composition of the alloy used in this investigation

is Fe25Cr7Ni3.5Mo0.25NWCu (wt.%). The present

study applied FSW using polycrystalline cubic boron nitride

(PCBN) tool to SAF 2507 super duplex stainless steel. The

postweld microstructure and mechanical properties of the FS

weld were examined.

2. Experimental procedures

Thebasematerialused in thepresent study is a commercial

SAF 2507 (UNS 2750) super duplex stainless steel, 4 mm

in thickness. Test plates were prepared with dimensions of

300 mm in length and 100 mm in width. Light sanding of

the top surface and joint butt surface to remove oxide and

contaminants prior to welding was performed using 80 grit

emery cloth. Before FSW, each coupon was degreased with

a methanol solvent.

FSW tool used in this study had a shoulder diameter of25 mm with the pin being 3.8 mm in length. The shoulder and

pin section of the tool were manufactured from solid PCBN.

A locking collar was used to hold the PCBN and transfer

torque from a tungsten carbide shank as shown inFig. 1.

FS welds were completedon a vertical milling machine fit-

ted with servomotors and control system. Since the machine

and tool were exposed to high temperatures during FSW, a

liquid cooled tool holder equipped with telemetry system to

broadcast tool temperature was used for the weld trails. An

argon atmosphere was introduced through a gas cup around

the tool at a flow rate of 2.8 105 mm3/s (1 m3/h) to avoid

the surface oxidation. A 3.5 tilt was applied to the tool dur-

ing FSW. The welding direction (WD) was identical to the

rolling direction (RD) of the plate. The welding parameters

were: rotational speed of 450 rpm and weld travel speed of

1 mm/s.

Microstructure in the weld was examined by opti-

cal microscopy (OM) and orientation imaging microscopy

(OIMTM)[26,27]. Sample for OM examination were elec-

trolytically etched in a 10 wt.% oxalic acid solution at 30 V

for 20 s. Details of the microstructure were examined by

OIM. Cross section for OIM was cut perpendicular to the

welding direction and then electrolytically polished in 20 ml

HClO4+ 180 ml C2H5OH solution at 223K (50C). Crys-

Fig. 1. PCBN friction stir welding tool assembly.

tallographic data were obtained from several regions, which

are illustrated as solid squares inFig. 2. The notation CEN

means the weld centre, which was defined as the centre of re-

gion swept by the shoulder. The notations ASn and RSn

indicate that the location of analysis isn mm away from the

weld centre at the advancing and retreating sides, respec-

tively. These notations are used throughout the present paper.

Crystallographic data collection by OIM was performed in a

PHILIPS XL30-SFEG scanning electron microscope (SEM),

operating at 30 kV under step size of 0.6m. Each observa-

tion area was 150m 150m. Crystallographic data were

expressed by phase map with grain boundaries and {1 1 1}pole figure. In the phase map, austenite and ferrite phases

were colored with gray and white, respectively, and the thick

and thin black lines show grain boundaries with misorienta-

tion exceeding 15 and misorientation between 3 and 15,

respectively. Coordinate axes of{1 1 1}pole figures are WD(or RD for the base material), transverse direction (TD) and

normal direction (ND) of the plate, as shown inFig. 2.

Vickers hardness test was conducted on the cross section

perpendicular to the welding direction, using a Vickers in-

denter with a 9.8 N load for 10 s, to examine the distribution


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378 Y.S. Sato et al. / Materials Science and Engineering A 397 (2005) 376384

Fig. 2. Schematic illustration of locations analyzed by OIM on cross section perpendicular to the welding direction.

of hardness originating with the duplex microstructure con-

sisting of ferrite and austenite. Transverse tensile specimens

were removed perpendicular to the welding direction and

prepared in accordance with ASTM E8. Tensile tests werecarried out at room temperature on a 445 kN MTS tensile

testing machine at a crosshead speed of 0.05 mm/s. A 51 mm

extensometer was used to determine the 0.2% offset yield

strength.

3. Results and discussion

3.1. Microstructure distribution of the weld

Low-magnification overview of friction stir welded SAF

2507 duplex stainless steel is presented inFig. 3.In the cross

section, the left- and right-hand sides of the weld center are

consistent with retreating and advancing sides of the rotating

Fig. 3. Cross section perpendicular to the welding direction of friction stir welded super duplex stainless steel 2507. An arrow shows the fracture location for

the transverse tensile test.

Fig. 4. Optical micrograph and phase map obtained by OIM of the base material.


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tool, respectively[2833]. The stir zone is seen around the

weld center. Theborder between the stir zone and the thermo-

mechanically affected zone (TMAZ) is very distinct on the

advancing side, while it is more diffuse on the retreating side.

It is apparent that the weld interior exhibits a high degree of

continuity and no defects.

An optical micrograph and a phase map of the as-receivedbase material are shown in Fig. 4.The base material has a

typical microstructure of wrought duplex stainless steels con-

sisting of ferrite matrix with austenite islands. OIM analysis

revealed that the austenite islands contained a higher number

of grain boundaries (mostly twin type boundaries) than the

ferrite. OIM analysis also revealed that the ferrite content was

about 51%. Averagegrain sizes of austenite andferrite phases

in the base material were about 4.3 and 5.1 m, respectively.

Optical microstructures of regions A, B, C and D

shown inFig. 3are indicated inFig. 5. Region B lies on the

weld centre, and region D is located on the border of the stir

zone and TMAZ. Regions A and C are located around 2 mm

away from the weld centre at the retreating and advancing

sides, respectively. Region B has the microstructure consist-

ingof ferrite matrixwith themore elongatedaustenite islands.

Austenite islands of region B look finer than those of the base

material. Region A has the similar microstructure to region

B, while region C seems to contain finer austenite islandsthan region B. Distribution of the austenite islands is finest in

the stir zone at the advancing side, as shown in micrograph

of region D. In this region, D, the austenite in the stir zone

exhibits an average grain size of 2.2m lying immediately

adjacent to elongated austenite islands in the TMAZ.

Phase maps of regions RS4.5, CEN and AS4.5 in the weld

areshown in Fig.6. All regions consist of a ferrite matrix with

the austenite islands similar to that of the base material. Dis-

tribution of the austenite islands in regions RS4.5 and CEN

is similar to that in the base material, but the austenite islands

contain more grain boundaries than the base material. Region

Fig. 5. Optical microstructures of regions A, B, C and D shown inFig. 2.


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Fig. 6. Phase maps of regions RS4.5, CEN and AS4.5 in the weld.

AS4.5 has more finely distributed austenite islands than thebase material, and theaustenite islands of this region lie along

the border of the stir zone/TMAZ. Again, the average grain

size of the austenite phase was finest in this region, averaging

2.2m. In all regions, the grain size of the austenite phase

was smaller than that of the ferrite phase.

Phase map of region AS5.5, located just outside the stir

zone, is shown in Fig. 7. This has the roughly same mor-

phology of austenite islands as the base material. However,

the ferrite matrix in this region contains a relatively higher

density of sub-boundaries than that in the base material. Av-

Fig. 7. Phase map of region AS5.5 in the weld.

erage grain size of theaustenite phase in this region was about4.0m, which is the roughly same as that of the base mate-

rial, while average grain size of the ferrite phase was about

4.0m, smaller than that of the base material.

Grain size profiles of the austenite and ferrite phases are

indicated inFig. 8. The austenite phase exhibits a smaller

grain size than the ferrite phase throughout the weld. In the

stir zone, grain sizes of austenite and ferrite phases decrease

from theretreating side towardsthe advancingside.The finest

austenite and ferrite grains are observed in the stir zone near

the TMAZ. Ferrite content profile across the stir zone in the

weld is shown in Fig. 9. Ferrite content varies between 50 and

60% across the weld and TMAZs, which is slightly higherthan that of the base material. The higher ferrite content in

the stir zone is attributed to the exposure to high tempera-

tures during FSW and the fact ferrite is more stable than the

austenite at the higher temperatures in the duplex stainless

steel[1].

{1 1 1}pole figures of the ferrite and austenite phases inseveral regions in the weld are presented inFig. 10. The base

Fig. 8. Grain size profiles of the austenite and ferrite phases in the weld.


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Fig. 9. Profile of ferrite content in the weld.

material has the weak{1 1 0}1 1 2 orientation in the austen-ite phase and the relatively strong {0 0 1}1 1 0orientation

in the ferrite phase, where{h k l} is the crystallographic planeperpendicular to the ND and u v w is the crystallographic

direction parallel to RD. Regions RS4.5, CEN and AS4.5

have different texture components, both in the austenite and

ferrite phases, from those observed in base material. These

changes in texture are indicative of the intense deformation

of material during FSW[29,30,32].In region AS5.5, on the

other hand, the texture component in the austenite phase is

roughly the same as that of the base material, while the fer-

rite phase exhibits a different texture component unlike that

of the base material.

3.2. Microstructural evolution during FSW

During FSW, the duplex stainless steel is heated up to

high temperatures created by frictional heating arising from

rotation of the welding tool. It is well known that the mi-

crostructure of duplex stainless steel changes to the fully

ferritic structure when the steel is exposed to temperatures

higher than about 1573 K (1300 C) [1]. Numerical calcu-

lation using the present chemical composition and the ther-

mochemical database [34] estimated that formation of the

fully ferritic microstructure required temperature exceedingabout 1600 K (1327 C) in the present steel. The formation of

the fully ferritic microstructure is followed by nucleation of

austenite phase at grain boundaries of the ferrite grains dur-

ing the weld cooling cycle. Therefore, the stir zone should

have the coarse ferrite matrix with intergranular austenite

phases, if it is heated up to temperatures higher than 1600K

(1327 C) during FSW. In the present study, however, several

microstructural analyses show that all regions in the weld

have the microstructures consisting of the ferrite matrix with

austenite islands, although the distribution and morphology

of the austenite islands areinhomogeneous in the weld. These

results indicate that the maximum temperature during FSW

did not exceed 1600 K (1327 C).During FSW, the material in the stir zone simultaneously

experiences large shear stresses along the pin tool surface, as

well as exposure to high temperatures [29,30,32]. A previ-

ous study[33]investigating material flow of FSW using an

in situ marker technique illustrated that material located at

the advancing side of the stir zone is displaced around the

entire circumference of the pin tool, ending up at roughly

the same location it began, displaced only slightly forward

of its original position. As a result, this report suggests that

material in the stir zone at the advancing side receives the

most severe deformation, which would tend to break up the

austenite islands producing a smaller grain size. This corre-sponds well with the microstructures reported herein, in that

the finest austenite grains are observed along the advancing

side of the stir zone.

Fig. 10. {1 1 1}pole figures of the austenite and ferrite phases in the base material (BM), regions RS4.5, CEN, AS4.5 and AS5.5 in the weld.


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Thegrainsizeprofile(Fig. 8) showedthat theaustenite and

ferrite grains in the stir zone were smaller than those in the

base material. Additionally, the phase maps (Fig. 6)showed

that both the austenite and ferrite phases in the stir zone did

not exhibit a heavily deformed microstructure, e.g. many low

angle grain boundaries. Both the grain size profile and phase

maps suggest that dynamic recrystallisation occurred bothin the austenite and ferrite phases during FSW. It is gener-

ally known that dynamic recrystallisation easily occurs in the

austenitic stainless steels, while ferritic steels hardly experi-

ence dynamic recrystallisation because the ferrite phase has

a high stacking fault energy[35].

In the case of the duplex stainless steels, however, de-

formation is localized in the ferrite matrix at high temper-

atures, because the ferrite phase is relatively weaker than

the austenite phase[36,37]. Consequently, the recrystallised

grains are often formed in ferrite phase more easily than

in austenite phase[3639]. Some research[36,37]suggests

that the recrystallised grains in the ferrite phase are formed

by continuous dynamic recrystallisation, which is charac-terized by strain-induced progressive rotation of subgrains

with little boundary migration. In the present study, the du-

plex stainless steel experienced plastic deformation by the

rotating tool at relatively high temperatures during FSW. As

such, it is likely that the ferrite matrix in the stir zone under-

goes continuous dynamic recrystallisation through the same

scenario.

On the other hand, the morphology of austenite islands

in the stir zone was much different from that of the base

material, as shown in Figs. 5 and 6. This suggests that the

austenite islands also experienced intense plastic strain dur-

ing FSW, leading to dynamic recrystallisation in the austenitephase.

After dynamic recrystallisation, the recrystallised grains

grow during the on-cooling thermal cycle [40,41].As men-

tioned above, since the ferrite phase is more likely to undergo

dynamic recrystallisation than the austenite phase, as a result

of the high temperature deformation of FSW, it is likely that

the recrystallised ferrite grains would grow early than those

of the austenite phase. This is likely the reason why the ferrite

phase exhibits a larger grain size than the austenite phase in

the stir zone.

Region AS5.5, located just outside the stir zone, had the

similar morphology of austenite islands to the as-received

base material. Grain size and texture of the austenite phase

in region AS5.5 were roughly the same as those in the base

material, but the ferrite phase had the different texture com-

ponents than that of the base material. This result suggests

that only the ferrite matrix underwent deformation and re-

crystallisation during FSW because the ferrite phase has a

lower flow stress at elevated temperature than the austen-

ite phase in the duplex stainless steel, as mentioned above.

It is generally known in Al alloys that dislocations intro-

duced into the TMAZ rearranged or migrate at tempera-

ture producing a recovered microstructure [4245]. How-

ever, the ferrite phase in region AS5.5 did not exhibit a de-

Fig. 11. Hardness profile across the stir zone in the weld.

formed microstructure. This result suggests that the ferrite

phase in this region may undergo sufficient plastic strains

to induce continuous dynamic recrystallisation as a resultof the significant localisation of deformation in the ferrite

phase.

3.3. Mechanical properties and effect of microstructure

on mechanical properties

A typical transverse hardness profile of a FSW in 2507

super duplex stainless steel is indicated inFig. 11.Given the

fact that the ferrite content in the stir zone is roughly uniform

(Fig. 9), the increase of hardness in the stir zone suggests that

the hardness profile is related to the grain sizes of ferrite and

austenite phases in the weld (Fig. 8).Transverse tensile properties of the weld are shown in

Fig. 12(a). The as-FSW 2507 super duplex exhibits roughly

the same 0.2% offset yield and ultimate tensile strengths as

the base material, with the exception of the elongation. To-

tal elongation to failure based on the standard 51 mm gauge

length was roughly 50% of the base material. However, given

the amount of reduction in area as observed in Fig. 12(b),

the actual percentage of ductility of the FSW specimens is

likely much higher than reported. All tensile failures oc-

curred roughly 7 mm from the weld centre at the retreat-

ing side, i.e. near the border of the stir zone and TMAZ,

as shown inFig. 12(b). This is consistent with the data pre-

sented inFig. 11,where the failures tend to move toward the

TMAZ as a result of the higher hardness, which is propor-

tional to strength in the metallic materials[9,46],of the stir

zone.

It should be pointed out that the tensile specimens used in

this study did not have uniform thickness (seeFig. 3)across

the weld. Typically, the thinnest section of a FSW is located

at the centerline of the weld, as a result of the tilt angle used

during welding. This is a reason why the transverse tensile

samples consistently fractured near the border of the stir zone

and TMAZ, rather than in the base material, which exhibited

the lowest hardness.


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Fig. 12. (a) Transverse tensile properties of the weld and (b) cross section of the fractured tensile sample.

4. Conclusions

The present study examined the microstructure and me-

chanical properties of FSW in SAF 2507 super duplex stain-

less steel. FSW using PCBN tool produced high-quality

welds in the super duplex stainless steel. FSW significantly

refined the ferrite and austenite phases through dynamic re-

crystallisation. The smaller ferrite and austenite grains cre-

ated increased hardness and strength in the stir zone. As a

result, weld transverse tensile failures consistently occurred

near theborder of thestir zone andTMAZ, exhibiting roughly

the same yield and ultimate tensile strengths as the base ma-

terial.

Acknowledgements

The authors are grateful to Mr. J.W. Pew and Mr. J.N.

Ostler for technical assistance and acknowledge Prof. K.Ishida and Dr. I. Ohnuma for the thermochemical calcula-

tion. They also wish to thank Prof. H. Kokawa, Dr. S.H.C.

Park and Dr. J.-Q. Su for their helpful discussion.

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Microstructure and mechanical properties of friction stir welded SAF

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