acom86_1+2 Engineering Properties of Duplex SS (2205, 2307) .pdf

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acom AVESTA CORROSION MANAGMENT Engineering Properties of Duplex Stainless Steels Duplex stainless steels, and in particular grade UNS S31803, are gaining importance as engineering material in various demanding industries with special emphasis on gas and oil applications. Because of their two-phase structure, duplex stainless steels are more complicated than pure one-phase stainless steels and their success- ful use is very much dependent on the knowledge of how they behave under various conditions. The reports published hereon aim at increasing the knowledge about metallurgical properties, fabrication properties, and applications and they were presented at the conference "Duplex Stainless Steels '86" in The Hague, October 26-28, 1986. "Influence of Nitrogen on Weldments in UNS S31803" by M Liljas and R Qvarfort Page 2-12 "Textures and Anisotropy in Duplex Stainless Steel SS 2377" by W B Hutchinson, U v Schlippenbach and J Jonson Page 13-17 "Applications and Uses of Duplex Stainless Steels" by J Olsson and S Nordin Page 18-23 All rights reserved. Comments and correspondence can be directed to Dr Sten Nordin, Avesta Projects AB, P.O. Box 557, S-651 09 Karlstad, Sweden. Tel. +46(0)54-10 27 70. Telex 66108 apab s. Telefax +46(0)54-18 82 54. N o 1-2 1986

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acom86_1+2 Engineering Properties of Duplex SS (2205, 2307)

Transcript of acom86_1+2 Engineering Properties of Duplex SS (2205, 2307) .pdf

Page 1: acom86_1+2  Engineering Properties of Duplex SS (2205, 2307) .pdf

acomAVESTA CORROSION MANAGMENT

Engineering Propertiesof

Duplex Stainless Steels

Duplex stainless steels, and in particular grade UNSS31803, are gaining importance as engineering materialin various demanding industries with special emphasison gas and oil applications. Because of their two-phasestructure, duplex stainless steels are more complicatedthan pure one-phase stainless steels and their success-ful use is very much dependent on the knowledge ofhow they behave under various conditions.

The reports published hereon aim at increasing theknowledge about metallurgical properties, fabricationproperties, and applications and they were presentedat the conference "Duplex Stainless Steels '86" inThe Hague, October 26-28, 1986.

"Influence of Nitrogen on Weldments in UNS S31803"

by M Liljas and R QvarfortPage 2-12

"Textures and Anisotropy in Duplex Stainless SteelSS 2377"

by W B Hutchinson, U v Schlippenbach and J JonsonPage 13-17

"Applications and Uses of Duplex Stainless Steels"

by J Olsson and S NordinPage 18-23

All rights reserved. Comments and correspondence can bedirected to Dr Sten Nordin, Avesta Projects AB,P.O. Box 557, S-651 09 Karlstad, Sweden. Tel. +46(0)54-10 27 70.Telex 66108 apab s. Telefax +46(0)54-18 82 54. No 1-2 1986

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Influence of Nitrogenon Weldmentsin UNS S31803

byMats Liljas and Rolf Qvarfort,

Research and Development, Avesta AB, S-77401 Avesta, Sweden

AbstractWelding has a profound influence on the structure ofduplex ferritic-austenitic stainless steels. It is wellestablished that both corrosion and mechanical prop-erties of duplex stainless steel weldments are stronglydependent on the ferrite-austenite balance. Forexample, resistance to pitting and intercrystalline cor-rosion as well as impact strength can be impaired byhigh ferrite levels.

Much attention has been paid to controlling the ferrite-austenite ratio in welds by restrictions of welding par-ameters. However, the wide limits for the alloyingelements according to UNS S31803 may result inlarge variations of microstructure in the weld zone.

Tests with laboratory heats show that the pitting resis-tance of both base metal and autogenous GTAW weldsis greatly improved by increasing the nitrogen contentin the steel. However, the weld metal has always a lowerresistance than the base metal.

A more austenitic composition of UNS S31803 withminimum 0.15% N has been tested. The base metal ofthis variant contains 5-10 % more austenite than themore conventional type with equal amounts of ferriteand austenite. After welding with normal practices theferrite content of weld metal or HAZ does not exceed65 %.

Properties of weldments including GTAW, PAW, SAWand SMAW are presented. Due to the relatively lowferrite contents high ductility and impact strengths areobtained.

Microprobe measurements of element partitioningbetween the two phases indicate that nitrogen has acontrolling effect on the austenite reformation in welds.

Pitting corrosion tests of weldments in ferric chlorideshow that the base metal (HAZ included) in the highnitrogen variant has a very high resistance while attacksin HAZ are quite common in heats with low nitrogen. Thewelding parameters as well as sample preparationappear to have a great effect on the results.

IntroductionAlthough produced for approximately 50 years duplexstainless steels have experienced increased interest inrecent years, due to the use in oil and gas industry. The

development has in general been towards higheralloyed steels which are more corrosion resistant andone trend has been to use the duplex materials inenvironments where pitting or crevice corrosion mightoccur.

As these types of localized corrosion normally are initi-ated at the weakest point on a steel surface the indi-vidual phases in the material should have similar pittingresistance. The parent material, solution treated atabout 1050°C, contains typically 50 % austenite and50 % ferrite with a higher content of austenite formerssuch as nickel, nitrogen and carbon in the austenite anda higher content of ferrite formers such as chromiumand molybdenum in the ferrite phase. Many express-ions as regards the alloy content have been proposedfor guidance concerning the pitting resistance. If thepitting resistance equivalent PRE = % Cr + 3.3 x% Mo + 16 x % N is used for both phases, it can bedemonstrated that a nitrogen content of 0.18 % is re-quired in a steel of the type 22Cr 5Ni 3Mo to obtainequal pitting resistance of the two phases (1). A positiveinfluence of nitrogen on pitting as well as intercrystallineand stress corrosion cracking resistance of the parentmaterial has been shown by many authors (1-5).

As the duplex stainless steels during welding undergo atransformation to ferrite followed by nucleation andgrowth of austenite, the welded joints have a differentmicrostructure than the parent material. Normally thecooling rates are not sufficiently low to obtain an equi-librium austenite content and, with the wide composi-tional limits in the standard of e.g. UNS S31803, veryhigh ferrite contents could be expected in the weldmetal if too ferritic a composition is selected.

Therefore, much attention has been paid to controllingthe ferrite/austenite balance in welds by restrictions onthe welding parameters (6, 7). However, welding withheat inputs normally used in manual welding (0.6-1.6kJ/mm), the arc energy has a minor influence on themicrostructure compared to the effect of composi-tion (8).

It is well established and understood that a high ferritelevel has an adverse effect on ductility and impacttoughness. As concerns pitting corrosion it is shown inmany works that the weld metal has a lower resistancethan the parent material (2, 9,10,11). A post weld heattreatment (PWHT) can improve the pitting resistance(11) and it has also been found that the use of a moreaustenitic filler metal by adding nickel or/and nitrogengives better pitting resistance (8, 12, 13). Both thePWHT which restores a higher austenite content anduse of a filler metal with increased nickel will result in a

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higher austenite content which of course also improvesthe impact properties. The austenite level required forobtaining acceptable pitting resistance has been dis-cussed (12, 14, 15), but there are naturally many otherfactors such as pitting index of the filler metal, presenceof carbide or nitride precipitates and surface roughnessthat will influence the result (13). The pitting attacks arereported to occur in the ferrite phase of weld metals inas welded condition (14) but in the austenite phase inthe PWHT condition (10). As the ferritic solidtficationresults in insignificant microsegregation (16), the parti-tioning of the elements during the solid state transform-ation must play an important role.

The purpose of this paper is to show the effect of nitro-gen on weld metal microstructures, mechanical prop-erties, and pitting corrosion resistance of a duplexstainless steel of the type UNS S31803.

ExperimentalMaterialsThe duplex stainless steels used in this investigationwere all of the type 22Cr5Ni3MoN (UNS S31803 orW.-Nr. 1.4462).

Laboratory heats with variations in nickel and nitrogenwere made into 30 kg ingots. Compositions are listed inTable 1. The ingots were hot forged and cold rolled to1.6 mm strips which were annealed for 5 min at 1030°C,water quenched and pickled.

Material was also obtained from normal productionheats with varying nitrogen contents. The compositionsare given in Table 2.

In the weld tests where filler metals were used moreaustenitic compositions than the base material wereselected. As seen in Table 3 the filler metals had nickelcontents ranging from 8 to 9.5 %.

Table 1: Composition of laboratory heats

Heat Elements, percent by weight

No C Si Mn P S Cr Ni Mo N

1 0.020 0.38 1.16 0.011 0.015 22.4 5.1 3.03 0.1272 0.024 0.37 1.18 0.011 0.013 21.8 5.1 3.02 0.1683 0.021 0.37 1.19 0.011 0.011 22.4 5.2 3.16 0.2144 0.024 0.35 1.38 0.011 0.015 21.9 5.1 3.03 0.217

5 0.024 0.37 1.13 0.011 0.013 23.6 5.4 2.98 0.0956 0.023 0.33 1.14 0.011 0.013 22.1 5.6 3.13 0.1307 0.027 0.34 1.15 0.011 0.015 22.0 5.6 3.04 0.1448 0.023 0.37 1.17 0.011 0.013 21.8 5.6 2.95 0.1499 0.022 0.34 1.15 0.011 0.014 21.8 5.5 3.08 0.181

10 0.025 0.37 1.23 0.012 0.015 22.2 5.6 3.03 0.236

11 0.026 0.38 1.25 0.011 0.014 21.6 5.9 2.81 0.12212 0.022 0.40 1.17 0.011 0.012 21.9 6.0 2.93 0.14113 0.022 0.37 1.14 0.011 0.012 22.2 6.0 3,02 0.15714 0.022 0.38 1.20 0.011 0.015 21.7 6.0 3.01 0.179

Table 2: Composition of parent materials, production heats

Code Thickness Elements, percent by weight

mm C Si Mn P S Cr Ni Mo N

I 3-12.5 0.021 0.40 1.50 0.024 0.001 22.03 5.71 3.16 0.179II 2-5 0.018 0.42 1.45 0.026 0.002 21.95 5.62 3.05 0.168III 3 0.020 0.40 1.52 0.022 0.001 21.94 5.53 2.98 0.150IV 6 0.014 0.42 1.41 0.025 0.001 21.95 5.66 3.07 0.128v 3 0.017 0.48 1.71 0.028 0.004 22.78 5.58 2.94 0.106

Table 3: Composition of filler materials

Code Type Diameter Elements, percent by weightmm C Si Mn P S Cr Ni Mo N

T1 Wire 1.2-3.0 0.018 0.49 1.68 0.018 0.003 22.57 8.04 2.99 0.150T2 " 3.0 0.016 0.45 1.50 0.018 0.003 22.41 7.94 2.99 0.140T3 " 1.2 0.011 0.48 1.61 0.016 0.003 22.50 8.00 2.95 0.130

T4 Rod 2.5 0.017 0.71 0.83 0.019 0.020 21.47 9.61 2.90 0.131T5 " 3.25 0.019 0.94 0.82 0.023 0.015 22.08 9.49 3,35 0.136

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Tests with laboratory heatsThe 1.6 mm strips were autogenously GTA welded(bead on plate) with arc energies of 0.05, 0.1 and 0.2kJ/mm.

Cross sections of the welds produced with an arcenergy of 0.2 kJ/mm were studied metallografically andthe ferrite contents were measured by point counting.For preparation see "Metallographic examination",next column.The critical pitting temperature (CPT) was determined inFeCI3 according to ASTM-G48-A on welds producedwith arc energies of 0.05 and 0.1 kJ/mm. Specimens insize 25 x 50 x t mm with the weld in the centreand parallel to the shorter side were cut out from thesamples. All surfaces of the corrosion specimens werewet ground to 400 mesh.

The corrosion tests were made at gradually increasingtemperatures in steps of 2.5°C. New duplicate speci-mens were used at each temperature. The CPT wasdefined as the lowest temperature where corrosionattacks could be seen on the surface at approx. 20 xmagnification. Since the attacks found in this investi-gation were very shallow, visible corrosion was oftennoticed without any significant weight loss of thespecimen. Attacks on ground edges were omitted.

Test with production heatsFor documentation of the weldment properties of UNSS31803 with high nitrogen content weld tests were per-formed on 2-12.5 mm plate with the compositions givenin Table 2. The welding procedures were selected tosimulate both longitudinal and girth welding of pipes.Also fabrication of seam welded pipes was tested.

Welding

The welding conditions for 2-12.5 mm plate and for thepipes are listed in Tables 4-6. The used joint geometriesare shown in Figure 1.

12.5 mm parent material: Weld A1 resembles the weldsin pipes 1300-1303. However, pipes 1300-1301 werewelded without filler in the first two passes whilepipes 1302-3 had filler addition in all passes. A2 simu-lates a girth weld with SMAW in the filler passes.

6 mm parent material: Weld B1U corresponds to pipes1705-6 with no filler addition while weld B1M corre-sponds to 1703-4 where filler metal was used in bothruns. B2 is a simulated girth weld utilizing GTAW andfiller metal.

2-5 mm parent material: The welding conditions fortests with 2-5 mm material of different heats are givenin Table 6.

Some of the welded plates and pipes were post weldheat treated. Weld A1 was annealed for 10 min at1080°C and water quenched while welds B1U and B1Mwere annealed for 10 min at 1050°C and air cooled.Pipes 1300,1302,1703 and 1705 were mill annealed at1080°C and quenched by water spray.

Mechanical testing

The welded plates were submitted to tensile and bend-ing tests across the welds.

Charpy V impact testing was performed at -40°C ofweld, heat affected zone and base metal. The notch wasoriented perpendicular to the surface giving crackpropagation in the weld direction. Three specimenswere tested at each location.

The tensile testing of the pipes was performed on longi-tudinal specimens either of the base metal or the weldmetal. Impact testing and bend testing of the pipeswere done in the same way as above.

Metallographic examination

Metallographic sections transverse to the weld werefirst etched electrolytically in 10% oxalic acid to revealpresence of precipitates and then etched in Murakami'setchant for staining the ferrite phase.

The ferrite content was measured by point counting ofthe austenite phase according to ASTM E562-83 at400 x magnification. A minimum of 30 fields weremeasured with a grid containing 16 points giving a maxi-mum standard deviation of ±10 % ferrite at the 50 %ferrite level. Also, the extended ferrite number (EFN)was measured using a Magne-Gage with counter-weights as described by Kotecki (17).

To relate the ferrite content to the annealing tempera-ture, steels I and IV were annealed for 10 minutesbetween 950 and 1350°C and water quenched. Theresulting ferrite levels were measured as above.

Electron Probe Micro Analysis (EPMA)

To assess the partitioning of major alloying elementsbetween ferrite and austenite in weldments EPMAwas performed by point counting. The specimenswere polished and slightly etched in Murakami'setchant prior to the analysis to facilitate positioningof the electron beam on the different phase regions.The nitrogen measurements were carried out accordingto a direct calibration method using a set of Fe-N alloys(0-0.435 % N) as references for the EPMA (18). Theinstrument was an ARL-SEMQ; 10 kV, 200 nA was usedgiving an accuracy of the nitrogen determinations ofabout 20 % and a limit of detection of about 0.02 % N.

Figure 1Weld preparationsa) Weld A1 and pipes 1300-3b) Weld A2

c) Weld B1, pipes 1701-4and 2-3 mm material

d) Welds B2 and F2b

Pitting corrosion tests

The test procedure described earlier was followed, withone exception: The edges of the specimen were groundwith 120 mesh abrasive paper, but the two plate sur-faces were not machined. The specimens were thenpickled to remove the weld oxide from the surfacesprior to test in the FeCl3-solution.

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Table 4: Welding conditions for 12.5 and 6 mm plate, heat I

Weld Thickness Weld run Process Filler Voltage Current Travel speed Arc energyno mm no heat dia V A mm/min kJ/mm

no mm

A1 12.5 1 PAW T1 1.2 27.5 200 150 2.22 GTAW - - 12 220 250 0.63-5 SAW T1 3.0 32 400 550 1.4

A2 12.5 1 GTAW T1 1.6 15.5 160 130 1.12 SMAW T4 2.5 23 67 64 1.43-5 SMAW T5 3.25 27 73 42 2.8

B1U 6 1 PAW - - 27.5 175 200 1.4

2 GTAW - - 12 220 250 0.6

B1M 6 1 PAW T1 1.2 27.5 200 150 2.22 GTAW T1 1.2 12 220 220 0.7

B2 6 1 GTAW T1 1.6 13 132 105 1.02 GTAW T1 1.6 13 123 64 1.53 GTAW T1 1.6 13 105 37 2.2

Table 5: Welding conditions for longitudinally welded pipes

Weld O.D.x t Weld run Process Filler Voltage Current Travel speed Arc energyno mm no heat dia V A mm/min kJ/mm

no mm

1300-1 273 X 12.5 1 PAW - - 27 210 200 1.72 GTAW - - 11 200 180 0.73-5 SAW T2 3.0 32 400 650 1.2

1302-3 273 X 12.5 1 PAW T3 1.2 27 210 200 1.72 GTAW T3 1.2 11 200 180 0.73-5 SAW T2 3.0 33 375 650 1.1

1703-4 168 X 6 1 PAW T3 1.2 28 210 260 1.4

2 GTAW T3 1.2 10 200 200 0.61705-6 168 X 6 1 PAW - - 26 190 300 1.0

2 GTAW - - 10 190 230 0.5

Table 6: Welding conditions for tests with 2-5 mm parent material

Weld BM Thick- Weld run Process Filler Voltage Current Travel Arc Type ofno heat ness no heat dia V A speed energy joint

no mm no mm mm/min kJ/mm

C1B I 3 1 GTAW - - 10 60 300 0.10 Bead on plateC1U I 3 1 " - - 13 168 200 0.66 ButtC1M I 3 1 " T1 1.2 14 170 200 0.71 "

D1U II 2 1 " 11 95 300 0.21 "

D1M II 2 1 " T1 1.2 12 110 300 0.26 "D2U II 2 1 " - - 12 62 90 0.50 "D2M II 2 1 " T1 1.6 13 60 90 0.52 "E1U II 3 1 " - - 12 115 200 0.41 "E1M II 3 1 " T1 1.2 13 130 200 0.50 "E2 II 3 1 " T1 1.6 12 65 60 0.78 V-80°

II 3 2 " T1 1.6 14 92 100 0.77 "F1 II 5 1 " T1 1.2 15 190 200 0.85 V-60°

II 5 2 " T1 1.2 15 165 200 0.74 "F2a II 5 1 " T1 1.6 12 70 55 0.92 V-80°

II 5 2 " T1 1.6 13 110 120 0.72 "II 5 3 " T1 2.4 15 130 105 1.11 "

F2b II 5 1 GTAW T1 1.6 15 105 130 0.73 see Fig. 1II 5 2 SMAW T5 3.25 30 100 260 0.69 "II 5 3 SMAW T5 3.25 30 110 260 0.76 "

K1B V 3 1 GTAW - - 10 60 300 0.10 Bead on plateK1U v 3 1 " - - 12 125 200 0.45 ButtK1M V 3 1 „ T1 1.2 14 155 200 0.65 "R1M III 3 1 " T1 1.6 12 70 94 0.54 Butt

R2M III 3 1 " T1 1.6 12 70 86 0.59 "

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ResultsTest of laboratory heats with differentnitrogen contentsThe CPT data obtained for base metal and autogenousGTAW weld metal are plotted versus nitrogen content inFigure 2. For both base metal and weld metal there isan improvement when the nitrogen content is raised.The different nickel contents are not influencing theCPT of the base metal while there appears to be a bene-ficial effect of nickel as regards the weld metal. Figure 2also shows that the autogenous weld metal has a5-10°C lower CPT than the base metal.

The ferrite contents of the weld metal are plottedagainst the nitrogen content in Figure 3. An increasednitrogen content resulted in a reduced ferrite contentas expected.

Figure 2Influence of nitrogen on CPT in FeCI3, laboratory heats.

Figure 3Ferrite content versus nitrogen concentration in autogenousGTAW welds. Arc energy 0.2 kJ/mm, laboratory heats.

Tests of production heats

Mechanical tests

The results of the mechanical tests are shown in Tables7-9. All tensile tests in Tables 7 and 8 across the weldresulted in fracture of the base metal and showed highelongations of both as welded (AW) and PWHT condi-tions i.e. no signs of weld defects, which was confirmedby the bend tests. The impact energies of the HAZ andthe weld metal (Tables 7 and 9) were in general some-what lower than that of the base metal. It appears thatthe GTAW welds had the highest impact value (weldB2). Weld A1, containing PAW and SAW also had a highimpact toughness. Welds B1U and B1M, produced byPAW with and without filler, gave equal impact values.The SMAW (A2) showed the lowest impact toughness.A PWHT with air cooling only gave a slight improvementwhereas a PWHT with water quenching resulted in alarge increase of the impact toughness.

The tensile tests of welded pipes (Table 8) showed thatweld metal and base metal had similar tensile prop-erties in both as welded and PWHT condition. Also herethe bending with a bend radius of 1.25 x t was per-formed without signs of cracks.

The impact tests of the welded pipes (Table 9) showedthat the welds produced by PAW without filler and SAW(1301) had a somewhat lower impact toughness thanthe one produced with filler (1303). PWHT resulted in agreat increase and the two types of welds had equalimpact strengths. The relatively low values in as weldedcondition for the fusion line were probably caused bycrack propagation into the weld metal.

The welds of pipes with 6 mm wall thickness showedhigh impact energies also when no filler addition wasused. In fact the weld metal was slightly tougher thanthe base metal. The PWHT only resulted in a small in-crease of the impact values.

Microstructures

a) Base metal

The ferrite contents versus annealing temperature forsteels I and IV are illustrated in Figure 4. Steel I had5-10 % lower ferrite content over the entire tempera-ture range 950-1350°C. Figure 4 also shows that steel IVis fully ferritic at 1350°C while steel I contains about90 % ferrite at that temperature.

Figure 4Ferrite content versus temperature for steels I and IV.

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Table 7: Results of mechanical testing of 12.5 and 6 mm plate

Weld Condition Thickness Tensile test, across weld Charpy V impact energy* -40°C, J Bend testmm Rp0.2 Rm A5 Base Fusion Weld Bend Bend

MPa MPa % metal zone metal radius angle

A1 AW 12.5 490 726 28 161 164 113 1.25 x t 180°PWHT " 538 748 39 - 273 227 1.25 x t "

A2 AW " 540 765 29 160 70 42 2 x t "

B1U AW 6 586 812 31 133 60 74 1.25 x t "PWHT " 552 807 35 125 66 78 " "

B1M AW " 582 816 29 133 61 77 " "PWHT " 535 818 33 126 72 82 " "

B2 AW " 584 822 30 - 130 142 " "

* For 6 mm thickness half size specimens were used:AW = As weldedPWHT = Post weld heat treated

Table 8: Mechanical tests of welded pipes

Pipe Condition Dimension Tensile test Bend testno O.D.x t Base metal Weld metal Across weld Bend Bend

mm Rp0.2 Rm A5 Rp0.2 Rm A5 metal Rm radius angleMPa MPa % MPa MPa % MPa

1300 PWHT 273X12.5 660 776 34 599 784 37 820 1.25 x t 180°

1301 AW " 722 828 29 758 840 26 876 " "

1302 PWHT " 624 779 33 671 793 37 838 " "

1303 AW " 713 812 30 822 859 31 871 " "

1703 PWHT 168X6 598 785 33 551 752 36 827 " "

1704 AW " 627 795 31 667 771 24 885 " "

1705 PWHT " 604 792 31 610 786 34 846 " "

1706 AW „ 646 825 30 674 807 28 889 " "

Table 9: Impact test of longitudinally welded pipes

Pipe Condi- Charpy V impact energy -40°C, Jno tion Weld Distance from fusion line Base

metal 0 mm 2 mm 5 mm metal

1300 PWHT 99 >147 >147 >147 >147

1301 AW 32 42 103 83 192

1302 PWHT 97 114 153 177 244

1303 AW 42 48 75 116 -

1703 PWHT* 107 91 89 93 95

1704 AW* 93 64 83 85 92

1705 PWHT* 110 94 92 88 93

1706 AW* 94 89 77 81 97

* Reduced size 10 X 5 mm cross section of specimen

b) Weld metal

The obtained ferrite levels in the weld metals and HAZare shown in Tables 10 and 11. The typical ferrite con-tent in the HAZ of the as welded condition is about 50%and in the PWHT condition about 40 %. Autogenousweld metals contained maximum 65 % ferrite. Afterannealing these welds contain about 40 % ferrite. Due toa more austenitic composition of the filler metals, de-posited weld metals had a range of 30-50 % ferritedepending on the degree of dilution from the basemetal.

Table 10: Ferrite contents in plate welds

Weld Condi- Thick-

Ferrite content, % EFNno tion ness HAZ Weld metal Weld metal

mm Face Root Face Root Face Root

A1 AW 12.5 57 63 51 65 78 110

A1 PWHT " 35 56 31 44 51 64

A2 AW " 47 50 30 40 41 49

B1U AW 6 45 52 61 64 96 86

B1U PWHT " 37 41 41 41 61 60

B1M AW " 51 51 56 52 78 82

B1M PWHT „ 41 41 35 34 60 58

B2 AW " 49 53 38 42 65 68

Table 11: Ferrite contents in pipe welds

Pipe Condi- Dimension Ferrite content, % EFNno tion O.D. x t Wel metal Weld metal

mm HAZ Face Root Face Root

1300 PWHT 273 x 12.5 32-36 25 34 44 58

1301 AW " 44 45 45 80 75

1302 PWHT " 30-33 29 30 41 65

1303 AW " 51 48 56 73 77

1703 PWHT 168 X 6 31-36 31 34 56 57

1704 AW " 48 53 56 81 72

1705 PWHT " 35-39 38 36 60 66

1706 AW " 50-51 61 63 98 84

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The extended ferrite numbers (EFN) of the weld metalsare also included in Tables 10-11. When correlating theobtained data the following relationship was found.

Ferrite content (%) = 4.5 + 0.59 x EFN.

Due to the limited width of the HAZ no measurement ofthe EFN could be executed.

Typical microstructures of the as welded condition areshown in Figures 5-6. In the ferrite phase small precipi-tates can be observed in optical microscope. Theseprecipitates appear only outside a certain distance fromthe austenite phase and where the distance betweenaustenitic areas is small no such precipitates are pre-sent.

Annealing resulted in complete dissolution of the pre-cipitates.

Figure 5Microstructure of Weld B 1 M. 6 mm plate PAW with filler metal.Arc energy 2.2 RJ/rnm, 56 % ferrite. 400 x

Figure 6Microstructure of root run (GTAW with filler) weld B1M.Arc energy 0.7 kJ/mm. 52 % ferrite. 400 x

Another observation is that the weld metal microstruc-ture varied with the distance from the surface. This isillustrated in Figures 7 and 8. Under certain conditionsthe surface structure was almost fully ferritic and con-tained precipitates while, under other conditions, thesurface had a border of austenite.

Electron Probe Micro Analysis (EPMA)Results of EPMA measurements in ferrite and austeniteof base and weld metal in two autogenously GTAwelded steels are shown in Table 12. The distributionbetween the two phases expressed as the partitioningcoefficient

content in ferriteP= ——————————

content in austenite

Figure 7Microstructure at the root surface of weld R2M. GTAW withno root gas. 400 x

Figure 8Microstructure at the root surface of weld R1M. GTAW with10 l/min Ar as root gas. 400 x

Table 12: Composition (percent by weight) of ferrite (F) and austenite (A) in weld and base metal of two welds

Sample Condition Area Ferrite Cr Ni Mo N

no % F A F A F A F A

C1U AW BM 40 23.3 20.7 4.7 7.3 4.03 2.66 0.02 0.29

AW WM 64 21.8 21.8 6.1 6.4 3.22 3.11 0.04 0.42

PWHT BM 40 23.6 20.9 4.6 7.3 3.81 2.52 0.06 0.30

PWHT WM 49 23.2 21.7 4.8 6.4 3.76 2.88 0.05 0.30

K1U AW BM 59 23.5 20.2 4.9 7.4 3.34 2.17 0.01 0.21

AW WM 89 22.6 22.2 6.0 6.3 2.95 2.73 0.04 0.45

PWHT BM 57 23.8 20.5 4.9 7.7 3.40 2.18 0.03 0.21

PWHT WM 56 24.0 21.3 5.0 7.1 3.43 2.40 0.04 0.23

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are shown for the both conditions in Figure 9. The parti-tioning for Cr, Ni and Mo in the base metal is in goodagreement with results from other investigations (19).The nitrogen content is 5 to 10 times higher in the aus-tenite than in the ferrite. However, due to the low nitro-gen level, there is a comparatively large scatter in thevalues for the ferrite phase. Assuming that the nitrogenconcentration of the ferrite in samples C1U and K1U is0.04 and 0.03 % respectively, the P-values as shown inFigure 9 are obtained. In as welded condition the con-tents of the substitutional elements were about thesame in austenite and ferrite and thus the P-values areclose to unity. Nitrogen on the other hand was stronglyconcentrated to the austenite. After annealing at1050°C the P-values in the weld metal were approach-ing those of the base metal.

No principal difference between the two heats with dif-ferent nitrogen contents could be observed. Due to thehigh ferrite content in weld K1U the nitrogen content inthe austenite was higher than in weld C1U althoughtotal nitrogen was lower in K1U.

Figure 9Partitioning of Cr, Ni, Mo and N in base metal and weld metalof welds K1U and C1U.

Figure 10CPT in FeCI3 versus ferrite content in weld metals. Weldsaccording to Table 13.

Pitting corrosion tests

All corrosion tests were made according to ASTM-G48-A on as welded specimens. As mentioned earlier,the microstructure at the weld surface was often differ-ent from that within the weld. This is the reason why theweld surfaces were not ground, but only pickled toremove weld oxide prior to test, although it will limit thepossibility to compare these results with other investi-gations. In most published investigations the surface ofthe weld metal has been removed by grinding.

The CPT-values measured for flat welded specimensand welded pipes and the corresponding ferrite contentare given in Table 13 and Figure 10. The CPT-value forthe base metal varied between 35°C and 40°C for differ-ent plate thicknesses of heat I. The CPT-values plottedversus the ferrite content in Figure 10 are the valuesmeasured at the sites where the corrosion attacksstarted i.e. root or top side respectively, as specified inTable 13.

Table 13 a, b: CPT values, location of attack and micro-structure at weld surface for the different welds

Table 13 a: Plate welds

Weld Condi- Microstucture* CPT Location*no tion at weld surface in of attack

Root Top FeCI3 at CPT

A1 AW 0.9FN 0.5F 25 R WA1 PWHT A 0.1 F 30 T WA2 AW 0.5F 0.5F 25 R WB1U AW 0.7FN 0.7FN 22.5 T WB1U PWHT 0.5F 0.5F 27.5 R T WB1M AW 0.5F 0.4F 30 T WB1M PWHT 0.5F 0.5F 25 T WB2 AW 0.5F 0.6F 22.5 R WC1B " - 0.9F 7.5 T WC1U " 0.9FN 0.4FN 15 R WC1M " 0.9F 0.4F 15 R W

D1U " 0.7FN 0.9FN 15 T WD1M " 0.9FN 0.4FN 15 R WD2U " 0.9FN 0.2F 22.5 R WD2M " 0.8FN 0.9FN 22.5 R T WE1U " 0.9FN 0.8FN 12.5 R WE1M " 0.7FN 0.8FN 15 T WE2 " 0.9FN 0.9FN 17.5 R W

F1 " 0.9FN 0.8FN 20 R T WF2a " 0.5F 0.8F 30 R WF2b " 0.3F 0.4F 27.5 T WK1B „ - FN 7.5 T W

K1U " FN FN 12.5 R T WK1M " 0.9FN 0.9FN 15 R T W

Weld A1-C1M: Heat IWeld D1U-K1B: Heat IIWeld K1U, K1M: Heat V

Table 13b: Pipe weldsWeld Condi- Microstructure* CPT Location*no tion at weld surface in of attack

Root Top FeCI3 at CPT

1300 PWHT 0.5F 0.5F <20 R W1301 AW 0.5FN 0.2F >17.5 -1302 PWHT 0.5F 0.5F <20 R W1303 AW 0.8FN 0.2F 22.5 R W1703 PWHT 0.5F 0.5F 25 T W1704 AW 0.9FN 0.4FN 22.5 R W1705 PWHT 0.5F 0.5F 30 R W1706 AW 0.7FN 0.8FN 20 R W

* Symbols: Microstructure: austenite (A), ferrite (F) andferrite with nitride precipitates (FN). 0.5F denotes a fractionof 50% ferrite. Location; weld metal (W), root (R) and topside (T) respectively.

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The CPT value is a measure of the material's ability toresist initiation of pitting corrosion. In order to measurethe effect of different parameters on the propagation,the weight loss of the specimens at temperaturesabove CPT was determined. Some examples are givenin Figures 11 and 12. The difference between the highand a low nitrogen materials is obvious, also when thelow nitrogen material has been welded with a filler metalwith nitrogen addition, Fig. 12. At a temperature of 10°Cover the CPT the weight loss of the high nitrogenmaterial steel was about one fifth of that of the lownitrogen material, although the CPT value was almostequal for the different welds, 12.5-15°C.

DiscussionEffect of nitrogen and nickel on pittingresistance and microstructure(laboratory heats)The tests with the laboratory heats showed that nitro-gen had a clearly beneficial influence on the pittingresistance of both base and weld metal, see Figure 2.A similar influence has been found for base metaland simulated heat affected zones (1). The effect ofnitrogen in the weld metal was at least as high as in thebase metal. However, the weld metal appeared to havea lower pitting resistance than the base metal. A certainpositive influence of nickel could also be observed forthe weld metal while nickel seemed to have no effect forthe base metal. Therefore, as nickel normally is con-sidered to have a negligible influence on the pittingresistance in stainless steels the observed positiveeffect should be related to the role of nickel in control-ling the weld microstructure.

The ferrite contents obtained in the weld metal with anarc energy of about 0.2 kJ/mm were in the range 70-90%. A multiple regression analysis on the influence ofchromium, nickel and nitrogen on ferrite content onlygave a statistically significant correlation for nitrogenhaving a clearly negative effect. As both the fraction ofaustenite and CPT of the weld metal were increased bynitrogen addition one conclusion could be that themajor contribution of nitrogen to increased pitting resis-tance is its effect on microstructure.

Test with production heatsMechanical properties

The mechanical tests of welded plates as well aswelded pipes showed that the strength and ductility ofbase and weld metal were similar both in as welded andPWHT condition. The impact tests were performed at-40°C, a temperature specified in many pipe-line pro-jects. Both GTAW and PAW processes gave high impacttoughnesses of weld metals in as welded and PWHTcondition also when no filler metal was used. The ferritecontents ranged from 30 to 60% but appeared to haveno significant influence on the impact toughness. Thiscould be concluded from Tables 9 and 11 where it isshown that the PWHT of pipes 1703 and 1705 has notresulted in any great improvement of the impact energyalthough the ferrite content decreased with more than20 %.

figure 11Weight loss in FeCI3 versus temperature for autogenous GTAwelds in 3 mm material.

Further data:

Weld %N % ferrite Arc CPT

energy

C1U 0.179 65 0.66 15D1U 0.168 70 0.21 15E1U 0.168 68 0.41 12.5K1U 0.106 81 0.45 12.5

Figure 12Weight loss in FeCI3 versus temperature for 3 mm GTA weldswith filler.

Further data:

Weld %N % ferrite Arc CPT

energy

C1M 0.179 60 0.71 15D1M 0.168 71 0.26 15E1M 0.168 58 0.50 15K1M 0.106 81 0.65 15

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The SMAW and SAW weld metals had impact tough-nesses of 30-40 J. Here, the PWHT resulted in a sub-stantial improvement of the impact energy with thesame decrease in ferrite level as for PAW weld metals.As the microstructure in the SAW weld metal was freefrom precipitates it is suggested that a coarser ferritegrain size and an unfavourable grain orientation causedthis difference in behaviour.

In general the impact energies noted for the fusion lineare on the same level as those of the weld metal. Due tothe weld geometry the crack propagation was partly inthe weld metal. Also the values 2 and 5 mm from thefusion line are in some cases influenced by weld metalwhich partly explains the lower values than for the basemetal.

Microstructure and composition

Autogenous GTAW and PAW weld metals of the highnitrogen steels, made with arc energies of 0.5-0.7kJ/mm contained 60-65 % ferrite. This could be com-pared with the heat with low nitrogen content whereferrite levels up to 90% was measured. However, whenvery low arc energies were used, as in the exampleswith GTAW bead on plate resulting in 0.1 kJ/mm, theferrite content was about 90% also in the high nitrogenmaterial.

When using low arc energies precipitation of fine par-ticles can occur within the ferrite in both weld metaland HAZ. Normally there is a precipitate free zoneimmediately adjacent to the austenite and in welds withsufficiently high austenite contents the distance be-tween the austenitic areas appears to be small enoughto prevent the precipitation. The precipitates were in alimited number of samples identified by electron diffrac-tion as Cr2N. In general a higher nitrogen content in thesteel appeared to reduce the occurrence of Cr2N byincreasing the amount of austenite. However, at low arcenergies resulting in ferrite contents above 60% somenitrides could be observed. Despite the presence ofnitrides in GTAW root welds the ductility and bendabilitywas very high.

The occurrence of a surface zone with higher ferritecontent than the bulk of the weld probably is the resultof different cooling conditions at the surface. Thisphenomenon has to be studied in more detail to be fullyunderstood.

When the weld metal microstructure is developed froma fairly homogeneous ferrite to a duplex ferritic-austen-itic structure, EPMA showed that only a very limitedmigration of substitutional elements is involved. On theother hand, nitrogen, being an interstitial with highdiffusivity, appears to participate to a very large extentin the reformation of austenite. The nitrogen levelsreported for the austenite in Table 12 were measuredon the coarser austenite areas along the ferrite grainboundaries. Limited measurements in smaller austeniteparticles within the ferrite grains indicated even highernitrogen contents. Thus, nitrogen, and most likelyalso carbon, plays an important role in controlling theaustenite formation. As a consequence, substitutingnickel with nitrogen could increase the austenite con-tent of the weld metal without altering the austenitecontent in the base metal.

Pitting corrosion

The results in this investigation show that the ferrite-austenite ratio in the weld surface is an important factorin determining pitting corrosion resistance, which tosome extent confirms the results in other investigations

(13). The metallographic examination of the structure atthe weld surface confirmed the detrimental effect ofvery high ferrite levels. In some cases extremely lowCPT values were correlated to very high ferrite levels,often with nitride precipitates also at the surface. Thespecimens with the highest CPT values had less ferritein the weld surface, sometimes even pure austenite,and no nitrides. The EPMA of the weld metal showedthat the alloy contents in the two phases were almostequal, the higher nitrogen content of the austenitebeing the only difference. This will have the effect thatthe PRE number (defined as PRE = Cr + 3.3 Mo + 16 N)for the ferrite is lower than for the austenite. Nitride pre-cipitates make the ferrite in the weld even more suscep-tible to pitting corrosion, and could therefore enhancethe effect of a low PRE number.

As the addition of nitrogen does not change the compo-sition of the ferrite in the weld metal to any large extentthe effect of nitrogen on pitting resistance should bethrough its influence on the phase balance.

The tests with production heats which were performedon pickled surfaces generally resulted in lower CPT thanthe tests with laboratory heats which were performedon ground surfaces. Apparently the weld surface has anunfavourable microstructure as regards the initiation ofpitting. The observed beneficial effect of nitrogen in thelaboratory heats was therefore not so evident on theproduction heats although a favourable effect of nitro-gen on the resistance to pit propagation at tempera-tures slightly above CPT was observed.

As all tests showed a lower CPT in the weld metal than inthe base metal (both with pickled and ground surfaces)filler metals overalloyed in chromium and molybdenumshould be considered for applications where a slightdecrease of the pitting resistance in the weld metalcannot be accepted. However, steels with low nitrogencontent often are attacked in the HAZ which would thenbe the weakest point. In the material with high nitrogencontent, on the other hand, no attacks have occurred inthe HAZ.

ConclusionsThe present investigation on duplex stainless steels oftype 22Cr 5Ni 3Mo N (UNS S31803) with variation innitrogen has shown that

- High nitrogen content gives high ductility and impacttoughness in the weld metal and heat affected zonein the as welded condition irrespective of if filler isused or not.

- The ferrite content after normal welding proceduresof the high nitrogen variant is 65 % maximum.

- Nitrogen is governing the reformation of austeniteduring welding.

- Nitrogen is beneficial as regards the pitting resis-tance of the base metal probably because i t in-creases the pitting index of the weakest phase whichis austenite.

- In the weld metal a positive effect of N can be ob-served on the bulk structure (ground surfaces) dueto the higher austenite content. However, the surfacestructure can vary considerably and high ferrite con-

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References1 Tsuge, H, Tarutani Y and Kudo T, Corrosion/86 Paper

No 156, NACE 19862 Miyuki H et al, Duplex Stainless Steels, ASM Metals/

Materials Technology Series, Paper 8201 -005,1983,p. 95

3 Sakai, J et al, J Materials for Energy Systems, Vol. 5No 2, Sept 1983, p. 105

4 Ishizawa Y, Corrosion/86, Paper No 159, NACE 19865 Guha P, dark CA, Duplex Stainless Steels, ASM

Metals/Materials Technology Series, Paper 8201-018 1983, p. 355

6 van Nassau, L, Welding in the World, Vol. 20, No 1/2,1982. p. 23

7 Mundt R and Hoffmeister H, Stahl u Eisen No 12,1983, p. 611

8 Gooch, T G, Duplex Stainless Steels, ASM Metals/Paper Materials Technology Series, Paper 8201-029, 1983, p. 573

9 Pleva, J and Nordin, S, Duplex Stainless Steels, ASMMetals/Materials Technology Series, Paper 8201-030, 1983, p. 603

10 Pleva, J and Johansson, B, Corrosion/84, Paper No218, NACE 1984

11 Kolts, J et al, Corrosion/82, Paper No 190, NACE1982

12 Pleva, J, Proceedings of "Stainless Steels '84", Göte-borg. The Institute of Metals, London 1985, p. 343

13 Bower, E N, Fielder, J W and King, K J, 25 Journeesdes Aciers Speciaux, St. Etienne, May 1986

14 Sridhar, N, Flasche, L H and Kolts, J, Corrosion/84,Paper No 244, NACE 1984

15 Ikeda A et al, Corrosion/86, Paper No 333, NACE1986

16 Suutala, N et al, Met Trans Vol. 10A, Aug. 1979,p.1183

17 Kotecki D. J., Welding Research Suppl. Nov. 1982,p. 352s.

18 Runnsjö, G, Carbon and Nitrogen in Steel ThesisUppsala University, 1981

19 Nagano H., Kowaka M., Tetsu-To-Hagane, 66 (8),July 1980, p. 1150

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Textures and Anisotropyin Duplex Stainless Steel

SS 2377by

W. B. Hutchinson, Swedish Institute for Metals Research, Stockholm, Sweden,U. v. Schlippenbach, RWTH, Aachen, FRG, and

J. Jonson, Research and Development, Avesta AB, Avesta, Sweden

AbstractCrystallographic textures have been measured in anumber of plate samples of duplex stainless steelSS 2377 and in one case the textures have been exam-ined in detail using orientation distribution functions(ODFs).

With decreasing plate thickness (i.e. increased hotrolling reduction) the textures become sharper and theduplex microstructure is progressively refined. As aresult of these changes the strength rises and becomesincreasingly anisotropic.

Calculations of mechanical anisotropy carried out usingthe ODF measurements show fair agreement withmeasured tensile test values. There is no evidence tosupport a significant fibre reinforcement effect arisingfrom the duplex microstructure.

IntroductionThe crystallographic texture and mechanical propertiesof duplex steel SS 2377 were reported in a recentpaper (1). It was shown that both the ferrite and aus-tenite were remarkably sharply textured and that thestrength of the rolled plate was unusually anisotropicas a result of the contributions from the two phasesand their mutual interaction. There was no evidence ofcomposite reinforcement effect controlling the be-haviour of the two-phase material. The present workextends that reported earlier to examine the situationin plates subjected to a range of hot rolling reductions.The semi-quantitative analysis of mechanical aniso-tropy which was presented earlier has been furtherrefined by the application of orientation distributionfunctions to the textures of the two phases.

Materials andexperimental procedureAll samples have been taken from commercially manu-factured plates of this alloy which has a typical com-position 22% Cr, 5.5% Ni, 3% Mo, 1.5% Mn, 0.5% Siand 0.15% N. The plates were produced by hot rollingwith subsequent annealing and quenching to equalisethe proportions of α and γ phase. Since the cast slabswere of constant thickness the accumulated strainduring hot working increased with decreasing final platethickness.

Mechanical properties have been assessed by stan-dard tensile testing of plates of various thicknesses.Microstructures have been examined by optical micro-scopy, and crystallographic textures have beenmeasured using reflection pole figures. Previous work(1) showed that the textures could vary considerablywith depth below the surface of the rolled plates. Thepresent measurements were restricted to the central50% of the plate where the structure and texture werealmost homogeneous.

The texture of the thinnest plate, D, was investigatedin greater detail using orientation distribution functions(ODFs) calculated according to the method of Bunge(2). The ferrite and austenite phases were treatedindependently; reflections from overlapping diffractionlines such as (110)α, (111)γ were omitted. Tensile testproperties were calculated from these ODF data on thebasis of Taylor's model (3) using the method describedby Hosford and Backofen (4) as elaborated by Bungeand Roberts (5). It was assumed in these calculationsthat the two phases were of equal volume and had thesame critical resolved shear stresses for slip. In previouswork (1) it was found that the microhardness of the twophases was quite similar. Results of the calculationswere compared with measurements carried out onspecimens from which 25% of the thickness had beenremoved from both sides by grinding in order toeliminate the heterogeneous texture, existing near tothe plate surfaces.

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Figure 1Representative yield stresses for tensile specimens fromrolling and transverse directions as function of plate thick-ness.

Figure 2Optical micrographs showing typical structures of the fourplates of different thickness. A = 30 mm, B = 15 mm, C = 8 mmand D = 4 mm.

Results and discussionFigure 1 shows a summary of tensile test data for platesof different thickness representing typical flat rolledproducts. It can be seen that the strength rises withreduction in thickness and that transverse samples areconsistently stronger than longitudinal ones.

This directionality of strength becomes more pro-nounced with increasing rolling reduction, being asgreat as 15% of the mean value for the thinnestplates. Such high degrees of anisotropy are seldomencountered in materials with cubic crystal structures.Optical micrographs in Figure 2 show the structures forfour plates of different thickness (A = 30 mm, B = 15mm, C = 8 mm and D = 4 mm). In all cases there isevidence of a laminated structure composed of alter-nating ferrite and austenite bands. However, the per-fection of the modulated structure becomes greaterwith increasing rolling reduction and the thickness ofthe individual lamellae decreases. Closer examinationshows that both phases are comprised of recrystallisedgrains. In almost all cases the ferrite and austenitelamellae are essentially one grain thick so the lamellaespacing may be considered to a good approximationas the grain size in these steels. There is evidently aprogressive refinement in grain size with decreasingplate thickness which is presumably responsible forthe increasing strength levels. To demonstrate thiseffect the yield strengths, taken as an average of the

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values in the 0° and 90° directions to minimise an-isotropic effects, have been plotted in Figure 3 as afunction of grain size.

Preferred crystallographic orientations of the plates Ato D are shown in Figure 4 in the form of (200) polefigures for the ferrite and austenite phases. Thetextures of the different samples contain broadly thesame components in each case but their sharpness ismarkedly enhanced with increasing rolling reduction.The thicker plates also show considerable asymmetrydue to residual effects from the coarse as-cast struc-ture. In the thinner plates the texture is remarkablysharp in both phases compared to single phase stain-

Figure 3Petch plot relating yield stress (average of 0° and 90° tests)to average grain thickness for the four plates A-D.

Figure 4(200) pole figures for ferrite (upper half) and austenite (lowerhalf) in the four plates A-D.

less steels and also differs in character. A completedescription of the texture in plate D has been preparedin the form of orientation distribution functions (ODFs)for both phases.

Figures 5 (a) and (b) show the ODFs for the ferrite andaustenite phases respectively. The ferrite texture con-sists principally of a sharp fibre spread from (100) [011]to (211) [011]. There are also minor spreads away fromthis latter component towards (311) [232] and (322)[241]. An interesting feature of the α-phase in this caseis the almost complete absence of (111) [uvw] com-ponents which usually dominate bcc rolling and anneal-ing textures. The austenite texture is well defined

Figure 5Orientation distributions for plate D (a) ferrite, and (b) aus-tenite.

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although not so sharp as that of the ferrite. It comprisesa peak at the orientation (110) [223] with ridges of highintensity connecting this to additional peaks at approxi-mately (230) [001] and further to (230) [324]. This is nota common preferred orientation in fcc metals but doesbear some similarity to that found after cold rolling andrecrystallisation of copper containing small additions ofphosphorus (6).

The textures observed in these duplex stainless steelsare difficult to interpet in that they may evolve bothduring hot rolling and during subsequent annealing.Temperature influences not only the tendency for re-crystallisation but also the relative proportions of thetwo phases. It is quite possible that the growth of onephase into the other with changing temperature is aform of recrystallisation which plays a role in deter-mining the final texture. Further work is necessary toelucidate these phenomena.

Using the ODF texture data shown above for steel D,the anisotropy of mechanical properties has been cal-culated for the ferrite and austenite respectively, andalso for the aggregate two-phase material. The calcula-tion (4,5) involves finding the most favourable condi-tions for plastic flow by varying the assumed strain ratio(r-value). The result gives the r-value and Taylorstrength factor which are predicted on the basis that allgrains undergo the same shape change. It is importantto note that calculations for the duplex case are notaverage values of the two phases separately. Mutualinteractions take place which can, for example, result inthe aggregate having a greater strength than both ofthe constituent phases.

Figure 6 shows the result of calculations of plastic strainratio (r-value) and Taylor strength factor (M) for tensiletesting along different directions in the plane of thesheet. The calculations were carried out assuming [111]<110> slip for the fcc γ phase and [110] <111> slip forthe bcc α phase. Some calculations were also madefor the latter phase based on <111> pencil glide. Forboth phases the textures produce a maximum in ther-value near to 45° to the rolling direction withminima at 0° and 90°. The anisotropy is considerablymore pronounced for ferrite than for austenite. The cal-culation for both phases together produces an inter-mediate behaviour which lies somewhat closer to ferritethan austenite. The Taylor factor which defines thestrength level is almost constant for austenite but variesgreatly in case of ferrite. There is a minimum near 45°and maxima at 0° and 90°. For the duplex material theTaylor factor is intermediate between the ferrite andaustenite values except close to the rolling directionwhere it lies above both of these.

Figure 7 compares the experimentally determinedr-values and flow stresses (at plastic strain = 1 %) withthe calculated values for the duplex material. The pre-dicted directionality of properties has the correct formalthough there are some substantial deviations from themeasured values. With regard to the r-values it is seenthat the predicted anisotropy is greater than that foundexperimentally although the reverse is true in the caseof the yield stresses. Note that the yield stress valueshave been fitted using the 0° test (rolling direction)since a calculation of the absolute strength level is notpossible.

It is often observed in bcc pure metals that slip is of apencil glide type rather than being restricted to [110]planes as assumed in the above calculations. Accord-ingly, some calculations were performed where theferrite phase was assumed to deform by <111> pencil

Figure 6Calculations of anisotropy of r-value (lower curves) and Taylorstrength factor (upper curves) using texture ODF data. Thecalculations are shown for both phases separately andcombined together in the duplex structure.

Figure 7Comparison of calculated and experimentally determinedanisotropy for duplex stainless steel D. Lower figure showsr-values and upper figure shows 1 % proof stresses.

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glide while the austenite retained the expected [111]<110> systems. These calculations for 0°, 45° and 90°test directions are shown by the stars in Figure 7.Somewhat worse fit was obtained regarding strengthlevels in this case than in the previous calculationsuggesting that pencil glide is not an appropriate de-scription for the slip mechanism in this type of highlyalloyed ferrite. The discrepancy between experimentand theory may to some extent be explained by theassumption of equal resolved shear stresses for slip inthe two phases, although microhardness values areobserved to be virtually the same in each.

There is no evidence in the present study of a fibrereinforcement effect contributing to the directionalityin strength properties. Such an effect would require asubstantial difference in strength between the twophases which is not observed, and should producemaximum and minimum strengths in the 0° and 90°directions respectively which is again contrary to ob-servations.

Finally, it should be noted that the calculations basedon texture data and the accompanying measurementsapply only to the central zone of the plate. In full-thick-ness tests the results are strongly influenced by thesubsurface zones which have differing textures andcontribute significantly to the difference in strengthlevels between 0° and 90° directions as shown inFigure 1.

References1 W.B. Hutchinson, K. Ushioda and G. Runnsjö, Mat.

Sci. & Techn. 1985, 1, 728.2 H.-J. Bunge, Mathematische Methoden der Textur-

analyse, 1969, Akademieverlag, Berlin.3 G.I. Taylor, J. Inst. Met, 1938, 62, 307.4 W. F. Hosford and W. A. Backofen, in Fundamentals of

Deformation Processing, Syracuse Univ. Press, 1964,p. 259.

5 H.-J. Bunge and W.T. Roberts, J. Appl. Cryst, 1961,2, 116.

6 V. Schmidt, K. Lücke and J. Pospiech, 4th Int. Conf.on Textures, 1975, Cambridge, p. 147.

ConclusionsThe strength of duplex stainless steel SS 2377 in-creases with refinement of the αlβ phase dispersionwhich may be considered as a grain size strengtheningeffect.

Both the austenite and ferrite may show unusuallysharp textures which result in a marked directionalityof strength. Fair agreement is obtained betweenmeasured anisotropy and that calculated on the basisof quantitative texture determinations. There is no evi-dence of a significant fibre reinforcement effect arisingfrom the duplex structure.

AcknowledgementsThe authors thank Professors R. Lagneborg and K.Lücke for provision of laboratory facilities. Useful con-tributions to the work were also made by K. Ushiodaand G. Runnsjö.

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Applications and Usesof Duplex Stainless Steels

byJan Olsson, Research and Development, Avesta AB, Avesta, Sweden, and

Sten Nordin, Avesta Projects AB, Karlstad, Sweden

Abstract Table 1: Typical chemical compositions of differentstainless steels mentioned in this paper

Duplex stainless steels offer a combination of strengthand corrosion resistance which makes them attractivefor a variety of applications. Today their major use is insour gas and sour oil. In this paper their engineeringproperties are discussed with emphasis on applicationsin chemical and petrochemical industries, environ-mental control, pulp and paper, food production andtransportation of chemicals.

IntroductionFerritic-austenitic or duplex stainless steels have beenavailable on the market for approximately 50 years.However, the very first duplex steels had an unfavour-able ferrite-austenite balance which resulted in ratherpoor mechanical properties of the heat affected zoneafter welding. Besides, the heat affected zone could beprone to intergranular corrosion due to a high carbonlevel in combination with the above mentioned un-favourable ferrite-austenite balance, viz. close to 100percent ferrite.

Today duplex steels are often composed in a waywhich ensures a high safety against the disadvantagesdescribed above and they should therefore be idealmaterials for general engineering purposes having thecombination of high mechanical strength, good cor-rosion resistance, good weldability and good formabilityin mind.

Despite this, the main field of application for modernduplex stainless steels have been piping for oil and gasexploration and to some extent heat exchanger tubing,which, from design and fabrication point of view, arerather simple applications.

The purpose of this paper is to emphasize the possibil-ities to utilize modern duplex stainless steels as generalpurpose engineering materials. The grades discussedare Avesta 3RE60, 25-5-1 L, 25-6-1 LN and above all,2205.

Historical backgroundOne of the first modern duplex stainless steels for gen-eral engineering purposes was Avesta 3RE60, intro-duced on the market close to 15 years ago. Comparedwith the traditional duplex stainless steels it had a lowerchromium level, which reduced the ferrite content. Inorder to maintain an acceptable corrosion resistancethe reduced chromium level was compensated for byan elevated content of molybdenum, see also Table 1.These measures gave a ferrite-austenite ratio of ap-proximately 60:40 compared to 70:30 for the oldgrades.

Grade Cmax Cr Ni Mo N Others

SS 2324* 0.10 25 5 1.5 -3RE60-old 0.030 18.5 4.7 2.8 -3RE60-modified 0.025 18.5 5 2.8 0.0825-5-1 L 0.025 25 5 1.5 0.0925-6-1 L 0.030 25 6 1.5 0.172205 0.030 22 5.5 3 0.15SS 2343 0.05 17 11 2.7 -304L 0.030 18 10 - -316L 0.030 17 11 2.3 -AISI 317 0.05 18 14 3.5 -904L 0.020 20 25 4.5 - Cu254 SMO 0.020 20 18 6 0.20 Cu

* Traditional duplex stainless steel

The first commercial installations of Avesta 3RE60 weremade late 1972 and the first failure due to intergranularcorrosion in the heat affected zone was reported in1973.

Two syrup tanks in a Swedish sugar plant, steel grade SS2343 (AISI 316), had suffered stress corrosion crackingin the bottom area and new lower parts, made of 3RE60,were installed early 1973. Late 1973 leakages werereported and at an inspection in January 1974 and bythe subsequent laboratory investigation it was estab-lished that the leakage was caused by intergranular cor-rosion. The heat affected zone contained almost 100percent ferrite and the carbon, which could not be dis-solved in the ferrite, had precipitated as chromium car-bides along the grain boundries.

The syrup contained up to 1 percent of chlorides, it hada pH of approximately 5 and it was stored at max 80°C.

A few months later similar damages were reported forhot water tanks and calorifiers in Australia. Most prob-ably the chloride containing water (up to 1000 ppmchlorides were reported) had initiated pitting and thenthe acid solution formed in the pits caused the propaga-tion of intergranular corrosion.

These cases clearly demonstrated that despite themore favourable austenite-ferrite ratio in 3RE60, thisfirst attempt to create a more versatile duplex stainlesssteel for general engineering purposes was not entirelysuccessful.

More austenite in the heat affected zone was required!

It should be noted, however, that the number of failures,(a few more were reported later) was very small com-pared to the total number of installations.

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Duplex of todayImprovement of 3RE60

The above described failures resulted in further devel-opment of 3RE60 aiming at a higher austenite content.

The parameter used was the ratio between the chro-mium and nickel equivalents acc. to the Schaeffler-deLong diagram and Table 2 shows a summary of the testresults leading to the modified 3RE60 with a high resis-tance to intergranular corrosion.

It was found that a material with a chromium equivalent- nickel equivalent ratio in the range 3.0-3.2 could besusceptible to intergranular cracking while a ratio of 2.6or less ensured sufficient resistance to this type of cor-rosion.

3RE60 has been produced according to this formula("modified" composition acc. to Table 1) since 1974without one single case of intergranular corrosionreported. More than 2000 tons of sheet and plate ma-terial have been supplied and the replacement of thesyrup tanks described above, see "Historical back-ground", was among the very first installations.

Engineering propertiesThe engineering properties of the duplex stainlesssteels are mainly determined by the ferritic-austeniticstructure and by the ferrite-austenite balance in theheat affected zone after welding.

Mechanical properties

The yield strength is approximately twice as high as forconventional austenitic grades and the tensile strengthis also higher. The ductility is inferior when compared tothe austenitic grades but still acceptable for severaloperations, see below under coldforming. The mostsevere deep-drawing operations, e.g. the pressing ofsheets with complicated patterns for plate heat ex-changers, cannot be conducted, however. The tough-ness is acceptable for sub-zero temperatures, also inwelded condition for temperatures down to at least-40°C.

Fatigue

The duplex stainless steels in general have far betterresistance to fatigue than the austenitic grades both ininert environments and corrosive environments. In thelatter case, the alloy composition and the corrosionresistance also play an important role, see Table 3.

Hot forming

The duplex grades are readily hot formed and at lowdeformation rates they show a superplastic perform-ance at temperatures in the range 800 to 1000°C.

Cold forming

All types of coldforming operations can be performedbut the lower ductility does not permit so severe press-ing operations as with the austenitic grades. Also theduplex grades workharden at coldforming but not to thesame extent as the austenitic grades. Spinning, e.g. forthe fabrication of dished ends, can be performed. Table4 shows some results achieved at the test spinning ofcones.

Table 2: Intergranular corrosion in neutral andacidified chloride containing waters.Welded 3RE60 U-bend samples.

Table 3: Fatigue and corrosion fatigue strengths inwater for some stainless steel grades.

Cl PH °C Cr equiv./Ni equiv.

ppm 3.2 3.0 2.6

5000 1 50 X 0 03 X 0 0

7 X - -

5000 1 70 X X 0

3 X 0 0

7 X 0 0

5000 1 90 - X 0

3 X X 0

7 X 0 0

10000 1 50 - 0 0

3 X 0 0

10000 1 70 X X 0

3 - 0 0

10000 1 90 - - 0

3 - 0 0

15000 1 70 X X 0

3 - 0 0

15000 3 90 X 0 0

Chlorides added as sodiumchloride, acidified by additionsof hydrochloric acid

x = intergranular crackingO = no intergranular cracking- = not tested

Grade Distilled water Synthetic sea water

SS 2343 (20) 225 MPa 215 MPa3RE60 (25) 320 MPa 305 MPa2205 (50) not tested 365 MPa

Figures within brackets give test material plate thickness.

Rotating bending with 1500 cycles/min. fatigue limits aregiven for 107 cycles.

Table 4: Results from the spinning of welded conesmade of austenitic and duplex stainlesssteels.

SS 2343 2205

Wall, mmoriginal 3.0 3.1spinning 1 2.5 2.7spinning 2 1.5 1.7

Cone length, mmoriginal 190 190spinning 1 215 210spinning 2 313 291

Final hardness, HvBase-matrix 327 380Base-weld 306 390Top-matrix 311 357Top-weld 345 413

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Welding

Welding can easily be performed utilizing the samewelding techniques as for austenitic grades. There aretwo hazards discussed in connection with the weldingof older types of duplex stainless steels

- too low heat input enables the heat affected zone tocool down too rapidly, thereby retaining an almostcompletely ferritic structure. This is in reality nohazard since the latest duplex steels, e.g. Avesta2205, have sufficiently high austenite stability also athigh temperatures.

- too high heat input causes grain growth of the ferriticgrains in the heat affected zone giving as a resultreduced toughness of this area. Several modernduplex steels e.g. Avesta 2205, have sufficiently highaustenite content of the heat affected zone to pre-vent such grain growth.

As indicated above neither low nor high heat inputimplies any great risk of deteriorating the materialproperties of the heat affected zone after welding, butusing a normal heat input means being on the safe side.

Corrosion resistanceOne consequence of the duplex structure is excellentresistance to stress corrosion cracking.

Resistance to general corrosion, pitting and crevicecorrosion is determined by the contents of alloyingelements, mainly chromium, molybdenum, and nitro-gen. When compared with the austenitic grades, Avesta3RE60 has approximately the same corrosion resis-tance as molybdenum alloyed steel of type SS 2343(AISI 316) while the higher alloyed grade 2205 issuperior to conventional grades like AISI 316 and 317and very close to Avesta 904L

Applications and usesAvesta 3RE60

Having in general the same corrosion resistance asSS 2343 with the extra merits of having superior resis-tance to stress corrosion cracking and corrosionfatigue, 3RE60 has mainly been used to resist SCC andcorrosion fatigue.

TanksAs mentioned above one of the first installations of3RE60 was for tanks where warm syrup was stored.After the first failure the 3RE60 tanks have been inservice for more than 10 years without any problemsreported.

Hot water tanks and hot liquor tanks for breweries areother examples where 3RE60 has been used success-fully for up to 7 years. The following data have beenreported for the stored liquids

chlorides up to 280 ppmpH 7.2-7.5temperature max 100°C

3RE60 has also been used for hot water tanks in othertypes of plants connected with food production, e.g.abattoirs.

Driers

One special type of drier has been used forthe drying ofspent grains from whisky distilleries and breweries,Fig 1. Rotating discs made of 3RE60 on a hollow shaftare internally heated with steam of 170°C. The spentgrains have a high moisture content, several thousandsppm of chlorides and a slightly acid pH.

Ten such "Rotadisc" driers made of 3RE60 are in servicein Great Britain and Ireland, the oldest since 1974.

Similar driers have also been used for the drying of fishmeal and blood meal.

Figure 1Rotadisc drier made of Avesta 3RE60 for a brewery.(Courtesy of Stord Bartz A/S)

Suction rolls

One of the biggest applications for 3RE60 has been suc-tion rolls for paper machines, Fig 2.

More than seventy such rolls have been supplied forpaper machines, the biggest being three rolls with alength of 10 m and a weight of 30 tonnes per piece for aUS paper mill.

Figure 2Suction roll made of Avesta 3RE60 for a paper machine.

Others

Avesta 3RE60 has also been used for a great number ofother applications, some of them listed below. In gen-eral, however, it has been used to solve stress cor-rosion cracking problems and in a few cases fatigue orcorrosion fatigue problems.

- Calorifiers- Centrifuges- Distillation columns- Expansion bellows- Fans- Heat exchangers- Process vessels

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Avesta 25-5-1 L and 25-6-1 LNLike 3RE60 the above mentioned grades are less com-mon today but the following examples show how duplexstainless steels can be used to solve specific corrosionproblems.

Pressure Vessel Fabrication

Four large pressure vessels for chemicals recovery in astyrene plant were fabricated. The process conditionsindicated risks for stress corrosion cracking and due tothe "non-productive" service the investment costs hadto be kept at a minimum. The engineering firm hadspecified the material to be 44LN duplex stainless steel,i.e. equal to Avesta 25-6-1 LN (R DirscherI and S Barth,"Duplex Stainless Steel Fabrication", 1982ASM MetalsCongress, Paper No 8201-026).

Design pressure was 2.7 atm (25 psig), max, and maxi-mum design temperature 260°C. At that time no ASMECode existed so the fabricator had to apply for a CodeCase to permit the material for welded construction inaccordance with Section VIII, Division 1 of the ASMEBoiler and Pressure Vessel Code. By the same token theweldability had to be evaluated and was found to begood and the results of the mechanical tests met therequirements of the ASME code as per above, SectionIX. The approval was designated Code Case 1893.

Two of the vessels were 4.5 m OD x 14.6 m with shellthickness 14 mm and head thickness 21 mm. The othertwo were 5.5 m ID x 11.3 m with thickness of 17 mm forshells and heads. Welding methods included SAW,SMAW and GTAW. After completed fabrication all buttwelds were fully radiographed. Only 8 m (out of 652 m)were rejected due to slag inclusions, 178 mm due toporosity and another 178 mm due to lack of fusion.

Not a single crack was found in a material that was be-lived to be susceptible to cracking.

This example shows that duplex stainless steels arewell suited for heavy vessels and that available fabrica-tion techniques can be used successfully.

Stripper Vessels in PVC Production

In the production of polyvinylchloride, PVC, whetherby the emulsion polymerization or the suspension poly-merization method, unreacted Vinyl Chloride Monomer(VCM) needs to be reclaimed and the PVC-slurry pu-rified from the monomer. In the VCM recovery stageVCM is recovered by simple flashing and stripping.However, in the stripping vessel, conditions that maycause stress corrosion cracking in austenitic stainlesssteels of the type 304L or 316L are at hand. In the midseventies, duplex stainless steel was supplied for suchvessels. For this particular process the operating tem-perature is about 85°C, pH value 3-3.5 and the waterphase has an average chloride content of 60 ppm, occa-sionally up to 150 ppm.

The fabricated vessels appear on Figure 3. They weremade from 5, 6,12 and 14 mm thick plates and weldingmethods were SMAW and SAW. The units went onstream in 1976 and have performed satisfactorily sincethen.

Refuse Incineration

A new concept for refuse incineration, utilizing the prin-ciple of wet air oxidation, has been developed in the US.Moist household refuse is being pressed down a verticalpipe several hundred meters into the ground. Pressur-ized air is added and the waste is being incinerated in anexothermic chemical reaction. Hence, the temperature

Figure 3Vessels for vinyl chloride monomer stripping made of Avesta25-5-1 L

increases down the tube and approaches 300°C beforethe waste has been fully oxidized.

The residual ashes are brought to surface in a jacketaround the pipe and then deposited. The high tempera-ture and the oxygen-rich atmosphere makes stress cor-rosion cracking a potential risk also at low chloride con-centrations. Accordingly, a duplex stainless steel wasselected for this service. The high strength simplifiedthe design, which had to take care of the weight of thepiping, and provide anti-collapse properties.

The system has operated for several years now and theduplex stainless steel has given excellent service.

Protein Recovery

Hydrocarbons are used for extraction of valuable in-gredients from various sources, e.g. waste from abat-toirs. For the extraction of proteins trichloroethylene isused.

From a corrosion point of view the use of this solventrequires special attention. In the presence of water ithas a tendency to decompose leaving free chloridesand the decomposition also brings about a decrease ofthe pH value. Hence, if austenitic stainless steels areused stress corrosion cracking may occur.

A Danish slaughter-house installed a system for therecovery of proteins with trichloroethylene some yearsago. The extraction takes place in a cylindrical vesselwith trichloroethylene flowing countercurrent to thefeed of meat remainders. Heating is done by steam sothe temperature is approximately 125°C. The shell ofthe extraction vessel as well as internals were made ofduplex stainless steel. It has behaved very well sinceput in service in 1980.

Avesta 2205

Contrary to the above described grades, 2205 has beenused not only to solve SCC-problems but also pitting,crevice corrosion and general corrosion problems.

FGD-Scrubbers

Close to 300 tons of 2205 sheet, plate and piping wasused for flue gas desulphurization scrubbers at a steelworks in New Zealand.

The gas enters a pre-cooler made of 904L, brick lined inthe hot parts, where it is cooled from 1000°C to 300°Cand then further to 150-200°C before it enters the ven-turi cooler scrubber. 2205 is used from the outlet of thepre-cooler and onwards. In the venturi the gas is cooledto 85°C and saturated with water, then it passes a fan

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Figure 4Precooler in a New Zealand FGD-unit made of Avesta 2205and 904L.(Courtesy of Fläkt Industri AB)

before it via a dewatering cyklone of 15 m height and 5m diameter is emitted through a stack, all made of 2205.

The recirculation liquor has a pH of 3-3.5, a chloridecontent of maximum 300 ppm and the temperature ismaximum 85°C.

The scrubbers were installed in 1984. Fig. 4 shows thepre-cooler made of 2205 and 904L.

Chimney in an FGD-system

Fig. 5 shows the fabrication of a 40 m high chimney in-tended for a Swedish pulp plant. The diameter is 4 m.

Flue gas from the soda recovery boiler is desulphurizedin a spray tower scrubber (made of Avesta 254 SMO)before emitted via the chimney. Despite the scrubbing,the emitted gas contains aggressive compounds likesulphur trioxide and hydrogen chloride which con-densed and caused corrosion on the previously usedchimney made of SS 2343.

The new chimney was installed in late 1985.

Figure 5A chimney made of Avesta 2205 under fabrication.

Pulp bleach plant filter washer

The environment in an acid stage pulp bleach plant filterwasher is always aggressive due to low pH, chloridesand residual chlorine from the bleaching process. Theconditions have been aggravated concurrently with theincreased demands from environmental authorities,forcing the plants to close their systems and recircu-late the liquids.

In one Swedish plant where SS 2343 previously hadbeen used for a filter drum it was decided to replace itwith a drum made of 2205 when they had to close theirwashing water systems.

The new 2205-drum was installed in the chlorine stagefilter washer in the summer 1985 and has performedwell during the first year of service. The following envi-ronmental data have been reported

temperature 35°Cresidual chlorine 12-15 ppmpH 2

Chloride level is not known but normal levels in C-stagewashers are several thousands of ppm.

Fig 6 shows the fabrication of this drum.

Figure 6Pulp bleach plant filter washer drum made of Avesta 2205under fabrication.(Courtesy of AB Hedemora Verkstäder)

Freon recovery units

In the tobacco industry freon is used to fluff up tobaccoleaves which have been packed for several years.Tobacco leaves are blasted with high pressure steamplus freon gas at 80°C. The freon takes up odours, con-taminants and steam. The spent freon is cleaned overactivated carbon beds in a column containing trays ofactivated carbon.

Previously AISI 304 was used for these units but pre-sence of chlorides caused pitting. Three units made of2205 were installed in early 1986.

Others

Avesta 2205 has also been used for several of the appli-cations described for 3RE60. Reason being because itis in general a more common grade with better avail-ability but also because it has been found that the cor-rosion resistance of 3RE60 in same cases has beeninadequate.

Some of these installations and also others are listedbelow

- Butadiene rubber holding tank- Calorifiers- Chemicals tankers, an entirely new field of applica-

tion for duplex steels where the combination of cor-rosion resistance and high strength (fatigue strengthincluded) are utilized.

- Fans- Flow lines for natural gas- Heat exchangers, tubular and fabricated from sheets,

for different types of industries, e.g. chemical andpetrochemical industries, pulp and paper plants.

- Hot water tanks and hot liquor tanks for breweries.- Press rolls for the pulp and paper industry (but not the

big suction rolls, yet).- Process reactors for brown coal.- Rotadisc driers- Valves

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acom is distributed free of charge to personsactively involved in the development of the processingindustry and other areas where stainless steels are im-portant.

acom appears four times a year, and we welcomeyour application, as well as additional applicationsfrom your friends and colleagues.

DiscussionThe purpose with this paper has been to show thatmodern duplex stainless steels can, fora great numberof applications, be used as conventional constructionmaterials where one or more of the following featuresare utilized,

- good corrosion resistance including other types ofcorrosion but stress corrosion cracking

- good fatigue and corrosion fatigue properties- good mechanical strength- good fabrication properties.

It is not necessary to consider this type of stainless steelonly when the environment implies a SCC-risk for aus-tenite grades. This is especially true for Avesta 2205,which has been used for applications where highlyalloyed and more expensive austenitic grades veryoften are selected.

to Dr. Sten Nordin, Avesta Projects AB, P.O. Box 557, S-651 09 KARLSTAD, Sweden

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