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WELDABILITY AND JOINTNABILITY OFDUPLEX STAINLESS STEELS
Prof. Dr. Srgio Duarte BrandiEscola Politcnica da Universidade de So
Paulo, Departamento de EngenhariaMetalrgica e de Materiais
Abstract. Welding and joining of stainlesssteels, in particular of duplex stainless steels
(DSS), has been one of the research lines of theWelding and Joining Group at Metallurgical andMaterials Engineering Department in University
of So Paulo since 1992. This paper is a briefreport on the work done in this subject up to2012. Weldability studies of multipass welding at
low temperature HAZ (LTHAZ) as well as at hightemperature HAZ (HTHAZ) and FZ (fusion zone)were conducted on samples of UNS S32101, S
32304, UNS S31803, UNS S 32550, UNSS32750 and UNS 32760 DSS. Microstructure,mechanical and corrosion properties in
simulated (dilatometer and or Gleebleequipment) and real welds were characterizedusing different techniques and tests.
Intermetallic phase precipitation (sigma phaseand chromium nitride) and secondary austenitewere characterized and a model for secondary
austenite precipitation in multipass welding ofHTHAZ was proposed. Simulated samples werecompared to real welds to validate the multipass
thermal model and presented a very good
agreement between them. Consequently thephase transformations studied by simulated
samples represent real weld microstructure. Thebehavior of multipass FZ using filler metals(AWS E2259-17 and EN 25 9 4L) was also
addressed. Intermetallic precipitation diagramswere determined for both filler metals and twoDSS to compare reheated weld metal regions
for three different welding procedures. Thecrescent heat input technique presented thebest results. Some dissimilar welding studies
using lean duplex, such as UNS S32101 andUNS S32304, and austenitic stainless steels
(AISI 304L and AISI 316L), were also carriedout. Brazeability of UNS S32101, S32304,S31803, S32750 and UNS S32707 DSS wasalso considered. Experiments were carried out
in a hydrogen continuous furnace using threedifferent nickel based filler metals (AWS BNi-1,BNi-2 and BNi-7). Several brazing conditions
were tested using different brazeability testssuch as the sessile droplet test, the edge test,and the capillary raise test. Best wetting and
spreading results were obtained for BNi-7 for theequipment used in the experiment, but brazing
thermal cycling in such furnace impaired basemetal corrosion resistance due to intermetallicphase precipitation, except UNS S32101 and
S32304.
Introduction
Weldability (and brazeability) of materials should
be analyzed as an interaction among materials(base material, filler metal, and weld metal),welding/joining process (thermal, mechanical,
and physico-chemical process characteristics),microstructure (FZ, PFZ, HAZ) and in-servicebehavior. Based in this approach, weldability
studies in DSS were accomplished using realwelds and simulated samples of fusion weldingprocesses [1-6].
To carry out theses studies DSS HAZ should be
divided in two parts, low temperature HAZ(LTHAZ) and high temperature HAZ
(HTHAZ)[2,3,7]. By low temperature one canunderstand a temperature range enough toavoid changes in the as -received microstructure
of DSS plates, which is from 950to 650
oC
depending upon DSS type. In other words,LTHAZ is a region of HAZ which has minor
changes in microstructure due to a small amountof intermetallic phases precipitation (chromiumnitride and sigma phase), usually not clearly
observed by optical microscopy, but enough tomodify mechanical and/or corrosion behavior[2].
On the other hand, HTHAZ is a region where as-received microstructure is completely modifiedand secondary austenite precipitation competeswith chromium nitride precipitation or dissolution,
depending on the peak temperatures andcooling rates of subsequent welding passesduring fabrication of an equipment. As a rule,
these temperatures range from solidustemperature to approximately 1000
oC.
Depending on the thermal cycle characteristics,
intermetallic phases and/or secondary austenitecan precipitate and might change locally theDSS properties.
Research results
HAZ research.
In the beginning, DSS weldability studies startedwelding UNS S31803 with autogenous one-passGTAW and EBW processes [1]. Heat inputs
were calculated to give a recommended t1200-800[8] for GTAW weld bead (approximately 9 s) and
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a lowert1200-800 for EBW (0,3 s). The amount ofaustenite observed in FZ and in HTHAZ of these
welds was (17 3) vol. % for GTAW and (3 1)vol. % in as-welded condition, whereas after 30min heat treatment at 1050
oC volumetric fraction
of austenite changed to (40 1) vol. % and (44
2) vol. % , respectively. Results of mechanicaland corrosion properties of post welding heat
treatment (PWHT) where similar to as-receivedplates. Figure 1 presents a result of generalizedcorrosion of EBW sample[9].
Fig. 1 Generalized corrosion in an EBW joint. (a) As-welded condition; (b) PWHT condition. Opticalmicroscopy. 40X. [9]
As PWHT is almost impractical for industrialconditions, the concept of a continuos PWHT
during welding come to light, and multipasswelding issue was addressed in further researchstudies. The idea involved in this concept was to
investigate the effect of multipass weldingthermal cycles in precipitation of intermetallicphases (sigma phase and chromium nitride);
and recovering the volumetric fraction ofaustenite (by nucleating intergranular and
intragranular secondary austenite in parallel withchromium nitride precipitation/dissolution) intomechanical and corrosion resistance of DSS.These phase transformations during welding
characterize LTHAZ and HTHAZ of a DSS,respectively.
LTHAZ research.
UNS S 31803, S32550[2], S32304, S32750,
and S32760[3] DSS samples were simulated in
a dilatometer and in a Gleeble machine usinga thermal modeling for LTHAZ[2,10] assuming
three passes at a place in the root joint, whichundergoes to a peak temperature of 950
oC in
the first pass, and using heat inputs from 0.4 to
1.0 kJ/mm. Intermetallic phase precipitation wascharacterized by extracting precipitates bydissolving the DSS matrix and carrying out X-
rays diffraction of extracted residues in a Debye-Scherrer chamber. Sigma phase and chromiumnitride precipitations were found in.
S32550[2,11], S32750, and S32760[3]. On theother hand, UNS S 31803 and S32304
simulated samples did not presented anyintermetallic precipitation detected by thistechnique. The amount of chromium nitride wascalculated for UNS S32550 and, it was
measured an amount of 0.35 vol. % for 1.0kJ/mm[2,11]. Thus, ductility of simulatedsamples of UNS S 31803 and S32550 where
compared using a bend test. Results arepresented in figure 2.
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Fig. 2 Bended surface of simulated samples with 1.0 kJ/mm welding heat input. (a) UNS S31803; (b)UNS S32550. SEM.[2]
Comparing figure 2(a) to figure 2(b) one cansee opened grain boundaries and interfaces insimulated sample of UNS S32550 due to
intermetallic precipitation at these regions, whichwas confirmed by ASTM A262 practice Aintergranular corrosion test. Also, pitting
corrosion resistance where measured usingartificial sea water and cyclic polarization tests attemperatures close to pitting potential critical
temperatures. Results are presented in table 1.
Table 1 Pitting and protection potentials for as-received and LTHAZ simulated samples of UNSS32304, S32750, and S32760. [3].
Material Condition Pitt ing potential(mV, SCE)
Protection potential(mV, SCE)
UNS S32304 (25oC test
temperature)
as-received 487 73 -117 41
simulated, 0.6 kJ/mm 531 37 -150 24simulated, 0.8 kJ/mm 459 32 -151 41
simulated, 1.0 kJ/mm 481 24 -164 10
UNS S32750(50
oC test
temperature)
as-received 1008 21 299 284
simulated, 0.6 kJ/mm 1073 25 -90 18
simulated, 0.8 kJ/mm 1110 10 -67 27
simulated, 1.0 kJ/mm 1083 11 -93 14
UNS S32760(50
oC test
temperature)
as-received 1030 37 409 138
simulated, 0.6 kJ/mm 1090 17 307 13
simulated, 0.8 kJ/mm 1053 25 250 29
simulated, 1.0 kJ/mm 1060 17 256 46
Analyzing table 1 one can note a significant dropin corrosion resistance observed for all heatinputs and materials studied compared to as-received condition, that is due to precipitation in
LTHAZ.In summary, all these results are regarding toLTHAZ, which is, as previously mentioned, a
region where almost no change in as-receivedmicrostructure is observed. Therefore, most of
the research published in literature in DSSwelding is related to HTHAZ, which presents ahuge change in microstructure in this region.
HTHAZ research.
During a stay at The Ohio State
University/Edison Welding Institute, a researchon simulated HAZ multipass welding of UNS
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S32550 and UNS S32750 was conducted usinga simple thermal cycle model to investigate
multipass welding effect on HAZ microstructureand mechanical and corrosion jointproperties[12,13]. Afterwards HTHAZ of UNS
S32205, S32304, S32550, S32750, and S32760were studied using GTAW welding process and
in a Gleeble equipment. A multipass thermal
cycle model for HTHAZ [14] was developed,using a distributed heat sources methodologyproposed by Grong and isten [15], with a
purpose of simulate HTHAZ samples.Fundamental aspects of intergranular andintragranular secondary austenite and chromium
nitride precipitation were considered and amodel of intragranular secondary austeniteprecipitation was proposed. In this model, a
cooperative precipitation of secondary austeniteand chromium nitride followed by nitridedissolution was proposed [4,16]. A scheme of
this model is presented in figure 3.Based on thismodel [4,16], chromium nitrides first nucleate
and precipitate at ferrite/austenite interface andgrowth into ferrite due to a favorable orientation
among ferrite/chromium nitride/austenite. Duringprecipitation, the regions adjacent to precipitatesare depleted in ferrite stabilizers alloying
elements, promoting austenite precipitation.Moreover, austenite has a higher nitrogensolubility and a dissolution of nitrides completely
involved by austenite takes place duringmultipass welding [16]. Mechanical andcorrosion properties were related with secondary
austenite precipitation. Charpy V absorbedenergy in multipass HTHAZ tends to be lowerthan as-received material[4]. Pitting corrosion
resistance was determined by artificial sea watercyclic polarization curves and results showed anincrease in pitting potential with reheating after
third pass and with increasing heat input.Results of pitting potential after third pass ofUNS S32550, S32750 and S32760 were almost
the same of as-received plates[4].
Fig. 3 Picture showing the dissolution of nitrides (white phase) in a secondary austenite region. Schemeof cooperative precipitation of secondary austenite and chromium nitride followed by nitridedissolution.[4,16]
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FZ research.
A similar research was conducted in fusion zoneof duplex stainless steels using two filler metals,
AWS E2209-17 and EN 25 9 4L. A TTPdiagrams were obtained for these two fillermetals and also to two DSS (UNS S32750 and
S32760). Comparing the results, filler metal EN25 9 4L presented the most favorable kinetic toprecipitate sigma and chi phases [5]. This
means that the filler metal is more susceptible tointermetallic precipitation than base metals
studied in this work, depending on weldingprocedure. Crescent heat input technique
presented better results than other testedwelding techniques. Fig. 4 shows the result forthe weld metal submitted to three reheating
thermal cycles.The effect of welding current frequency on theweld metal grain size was also studied for UNS
S32101 and UNS S32304, which is depicted infigure 5. In this figure is also shown, for UNSS32102, EBSD images to compare grain size in
a non-pulsed current and 20 Hz currentfrequency.
Fig. 4 Picture showing the TTT sigma phase precipitation curve for weld metal from EN 25 9 4L andthee reheating thermal cycles and the microstructure of the weld metal in this region.[5]
Analyzing figure 5 on can see a trend to
increase the weld metal grain size of UNSS32101. On the other hand UNS S32304presented also a slightly trend to increase the
grain size with the increase of the weldingcurrent frequency.
To investigate these results, the nitrogen contentwas determined for each weld metal, for bothbase metals. The results are presented in table2.
The loss of nitrogen was calculated based upon
the original amount of the base metal beforewelding.The nitrogen loss increased with welding current
frequency for both duplex stainless steels. UNSo nitrogen.S32101 presented a higher nitrogen
loss than UNS S32304. This behavior might bedue to the different chemical composition. UNSS32304 has a higher amount of chromium,which has a higher chemical affinity
AWS E 2610-17
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Fig. 5 Effect of GTA welding current frequency on the grain size of UNS S 32101 and UNS S32304. In
this figure is shown EBSD images of weld metal grain size for non-pulsed welding current and 20Hz
welding current frequency. [18]
Table 2 Nitrogen amount and nitrogen loss related to duplex stainless steel grades and differentwelding current frequency [18].
UNS S32101 Nitrogen (%)Nitrogen
lossUNS S32304 Nitrogen(%)
Nitrogen
loss
Base metal 0,213 - Base metal 0,120 -
no pulsed 0,153 0,060 no pulsed 0,080 0,040
1 Hz 0,125 0,088 1 Hz 0,065 0,055
5 Hz 0,110 0,103 5 Hz 0,064 0,056
10 Hz 0,091 0,122 10 Hz 0,060 0,060
20 Hz 0,110 0,103 20 Hz 0,047 0,073
A multicomponent phase diagram for UNSS32101 was built up to confirm the effect of
nitrogen in weld metal grain growth, aspresented in figure 6.
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As the frequency increases, the amount ofnitrogen decreases and, as a consequence, the.
temperature interval in the ferritic field increases.As the heat input is constant for all experiments,
the cooling rate is almost the same. It means,the higher the interval, more time in the ferritic
field, producing a larger grain size.
Fig. 6 - Phase diagram of the LDSS UNS S32101, a) Complete phase diagram, b) Enlarged area of the
ferritic field and the nitrogen loss according to welding current frequency, compared to base metal. [18]
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Dissimilar welding research
In some applications duplex stainless steels arewelded to other steels grades, such as carbonsteels or austenitic stainless steels. In figure 7 is
shown a microstructure of a similar welded jointof UNS S32304 and a dissimilar welded joint ofUNS S3234 and AISI304. These materials were
welded with coated electrode E2209-17.In figure 7 is also presented the result of ASTMA262 Practice A corrosion test in the regions
indicated. The results indicates a sensitizedregion in the HAZ of UNS S32304 and in HAZ of
AISI 304.
Fig. 7 A similar and a dissimilar welded joint of UNS S32304 to AISI 304. In the picture is shown theresult of ASTM A262 Practice A corrosion test. [19]
AISI 304 UNS S32304E2209-17
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Brazeability research.
Brazeability is a property which involves aspectsrelated to filler metal, base metal and itsinteraction with flux or atmosphere and thermal
effects in base metal during brazing [6]. In otherwords, filler metal, base metal and brazingprocess characteristics should be chosen in
such a way that no embrittlement of base metaloccurred during brazing, liquid metal flowthrough all the joint gap and joint region
protected by flux or suitable atmosphere. Basedon this definition, brazing temperature shouldavoid Change in volumetric fraction indicates
that other phases might precipitate dependingon temperature and time during brazing thermalcycle. For UNS S31803 the temperature range
to be avoided is from 950 to 700oC, due to
sigma phase precipitation in base metal. The
best brazing temperature for duplex stainlesssteels is around the heat treatment at 1050
oC.
Temperatures higher than this can produce achange in ferrite/austenite balance and might
produce nitride precipitation. For this reason,brazing thermal cycle should be carefullydesigned to avoid embrittlement in the whole
component. As duplex stainless steels have agood corrosion resistance, the brazing fillermetal should have also a good corrosion
resistance. Nickel base filler metals present
good corrosion resistance and a suitable brazingtemperature range. Nickel base filler metals are
alloyed with boron and silicon (BNi-2) andphosphorus (BNi-7) to reduce brazingtemperature [6]. These elements also affect
liquid viscosity and, as a consequence, gap jointfilling extension. Each filler metal presents anideal clearance to provide a defect free joint. All
these aspects together characterize brazeabilityof a material. In other words, contact angle,spreading final area, ideal joint clearance
determined by edge test, and phasetransformation in base metal during brazingthermal cycle are important to characterize
brazeability.
Brazing of UNS S31803 DSS was done in acontinuous furnace under a pure hydrogen
atmosphere and using different brazing fillermetals, BNi-2 and BNi-7 at 1100
oC. Brazing
parameters were changed to produce differentbrazing and after-brazing cooling times. Resultsshowed that BNi-7 presented better wettabilityresults than BNi-2 [6]. However the cooling time
after brazing was low enough to deteriorate DSSbase metal, producing approximately 10% ofsigma phase. Based on this result, continuous
furnace brazing are not suitable for UNS S31803DSS brazing, unless a cooling system might beused to reduce cooling time and avoid sigma
phase precipitation[6].
Fig. 8 Relationship between joint strength and joint gap [6].
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Edge test provides important informationregarding to ideal brazing joint clearance. The
minimum joint clearance that produces a jointcompleted filled and the maximum gap whereeutectic phase in the center of joint is
discontinuous define this ideal range. In thecase of duplex stainless steels there is a fccnickel rich phase adjacent to base metal and a
eutectic with hard and brittle phases in themiddle of the joint. There is a strong correlationbetween the microstructure and the mechanical
properties of a brazed joint. The mechanicalstrength of a brazed joint is achieved when thegap is into the ideal gap range. Figure 2
presents an schematic plot of the strength of thejoint as a function of the joint gap. Region 1presents a drop in the mechanical properties
due to a lack of joint filling, reducing the brazedarea in the joint. Region 2 shows the higher
strength, due to the Ni rich continuous
phase in the gap and also due to a mechanicalconstraint produced by a triaxial stress state.
Region 3 depicts a region with the eutecticphase in the joint middle, producing amechanical resistance close to the filler metal.
Figure 9 presents a microstructure of a non idealgap of a UNS S31803 brazed joint with BNi-2filler metal. In figure 3(a) one can see the rich
nickel phase at brazed joint/base metal interfaceand some rich in silicon and/or boron phases,represented by the dark grey phases. Figure
3(b) shows a precipitation of intermetallicphases in base metal. These phases are sigmaphase and boron rich phases. During brazing,
boron can diffuse to base metal faster thansilicon and phosphorus.
(a)
(b)
Fig. 9 Brazed joint of UNS S31803 with BNi-2. (a) shows the microstructure of the brazed joint. (b)depicts the intermetallic phases produced during brazing in base metal. Back scattered electronspictures.
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Figure 10 shows a microstructure of a brazedjoint of a UNS S31803 joined with BNi -7 filler
metal. In figure 4(a) one can see the rich nickelphase at brazed joint/base metal interface andsome phosphorus rich phases. Figure 4(b)
depicts a back scattered electrons picture of thejoint. The light gray in the joint was identified by
microanalysis as (Ni,Cr,Fe)P and the dark grayas (Ni,Cr,Fe)3P. Sigma phase precipitation was
also observed in these brazed samples.
(a)
(b)
Fig. 10 Brazed joint of UNS S31803 with BNi-7. (a) shows the microstructure of the brazed joint. Etch:10% oxalic acid. (b) depicts the intermetallic phases produced in the brazed gap. Back scatteredelectrons picture.
A EBSD technique was utilized to characterizeferrite and austenite in UNS S32750 joined with
BNi-2 brazing filler metal. Results are presentedin figure 11.
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Figure 11(a) shows the presence of a largeamount of nickel rich phase in the eutectic
phase, as previously presented.
(a)
(b)
Fig. 11 EBSD image of ferrite and austenite in the brazed joint of UNS S32750 with BNi-2. (a) shows
austenite and ferrite as solid colors. (b) depicts the region characterized by EBSD technique. Backscattered electrons picture.
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Final comments
This paper is an overview of the research done atUniversity of So Paulo in weldability and jointability
of DSS in the last 10 years. Details of tests and
additional results and discussions can be found inthe mentioned literature.
Acknowledgments
Author would like to acknowledge his students
(Antonio, Claudia, Ricardo, Clvis, Vinicius,Francisco, Adriano, Marcos, Reginaldo, Alcio, Jos
Antonio, Silveli),which make these ideas come
through. Also I appreciate FAPESP and CAPES forfunding these works.
References
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[15] O. Grong Metallurgical modeling of welding
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[19] Assis, C. A. - Comparao da junta similar de
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