acom85_4 high temperature behaviour of the austenitic ss uns s30815 (253 ma) & weldments .pdf

7/29/2019 acom85_4 High Temperature Behaviour of the Austenitic SS UNS S30815 (253 MA) & Weldments .pdf

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acomAVESTA CORROSION MANAGMENT

High TemperatureBehaviour of the Austenitic

Stainless SteelASTM UNS S30815

(253 MA) and Weldmentsby

A. Dhooge, Head of the Welding Research Dept, Research Center of the Belgian Welding Institute, Gent, Belgium,W. Hoek, Materials Engineer at DOW CHEMICAL (Nederland) B.V., Terneuzen, The Netherlands,

W. Provost, Research Engineer at the WTCM-CRIF, Gent, Belgium, andM. Steen, Research Engineer at the Laboratory Soete for Strength of Materials and Welding Technology,

University Gent, Belgium

IntroductionMaterials used in high temperature installations aresubjected to severe loads, such as creep, mechanicalor thermal fatigue, thermal shock, environmental attackand their interaction. In order to obtain a safe andreliable operation, the selection of the base material,weld metal and welding procedures is very importantand must be based on a complete knowledge of operat-ing environment, loads, temperature and externalloading conditions together with sufficient test data ofthe potential materials.

The weldments are the most critical areas, as they areobvious sites for crack initiation and propagation.A correct selection of welding procedures and con-sumables is therefore vital for the construction to reachthe design lifetime. In general, sufficient test data areavailable on base metals, but there is a lack of informa-tion about the behaviour of weldments at elevatedtemperatures.

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. No 4-1985


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In this paper some high temperature data on the fullyaustenitic Cr-Ni-N-RE material 253 MA* (ASTM UNSS30815) and weldments are presented. According tothe data of high temperature material vendors (1),253 MA shows some interesting advantages over thecommon high temperature materials AISI type 304H(UNS S30409) and Alloy 800H (UNS N08810):

- higher creep-strength than AISI 304H (Figure 1)- lower coefficient of thermal expansion than AISI type

304H (Figure 2)

- only about 60% more expensive than AISI type 304Hwhile Alloy 800H is about 2.5 times more expensive.

In the present paper, no attention is paid to the in-fluence, of the environment on the materials' high tem-perature behaviour, although it is recognized that suchinfluences may be very important. Attention is given tothe weldability and high temperature tensile and creepproperties of the base material and manual metal arc(MMA) weldments.

The hot cracking sensitivity is determined on the basis

of "hot ductility" tests, while the creep properties areevaluated using a newly developed testing methodbased on low constant strain rate tensile tests (2).Finally, the metallurgical stability is evaluated bymechanical testing and microscopic examinations afterlaborator a ein .

(*) 253 MA: tradename of Avesta/Sandvik

Test material

Figure 1Elevated temperature properties of three austenitic high tem-perature materials (1)

Figure 2Thermal expansion coefficient of three austenitic high tem-perature materials (1)

The test specimens are taken out of a 12 mm thickfully austenitic Cr-Ni stainless steel plate 253 MA; itschemical composition is given in Table 1. In addition tothe conventional alloying elements, this grade containsrare-earth-metals, added in very small but carefully con-trolled amounts, and has a high silicon content in order

to provide good oxidation resistance and high creepproperties. Some nitrogen is added to make the steelfully austenitic and which, together with the compara-tively high carbon content, provides good high tem-perature strength.

For comparison some specimens of Alloy 800H aretaken out of plate material of 15 mm thickness; thechemical composition is also given in Table 1.

Table 1: Chemical composition of high temperature materials under investigation.

%C % Mn % Si % S % P % Cr % Ni % AI % N Others

253 MA 0.090 0.56 1.52 0.001 0.024 20.80 11.20 - 0.179 0.05 Ce

Alloy 800H 0.084 1.05 0.37 0.002 0.004 19.28 31.13 0.23 - 0.34 Ti


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Hot cracking sensitivityOne of the problems encountered in welding hightemperature materials is the crack susceptibility in theheat affected zone (HAZ). The cracks are predomi-nantly hot cracks and are caused by low melting

eutectics which spread across the grain boundaries.The stresses due to welding are then responsible forthe occurrence of cracks.

The hot cracking susceptibility of the 253 MA basemetal was determined by performing "hot ductilitytests" (3). In this test, specimens are subjected to aweld thermal cycle and ruptured by high strain ratetensile loading at different temperatures on heating oron cooling (Figure 3).

In this way the temperatures at which respectively theductility and the strength drop to zero are determined.Once these so called nil-ductility (TND) and nil-strength(TNS) temperatures are defined, specimens are heatedto a peak temperature equal to the nil strength tem-

perature and subsequently pulled to fracture on cool-ing. The reduction in area and the ultimate tensilestrength are determined and plotted versus the testtemperature (Figure 4). For materials which are proneto the formation of hot cracks, a marked reduction intemperature occurs before an acceptable ductility isregained (TNS - T*ND >100C).

The test results obtained on 253 MA base metal(Figure 4) reveal that this metal is not susceptible toheat affected zone hot cracking. Indeed, both the duc-tility and the tensile strength recover rapidly whentesting specimens on cooling from the nil-strength tem-perature (TNS - T*ND = 30C).

Figure 4

Hot ductility results on cooling from TNS - 253 MA basematerial

Figure 3Hot ductility test


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Welding procedures12 mm thick plates are manually metal arc welded withthree different types of coated electrodes: one fillermetal is a matching electrode (weldment A), the othersa 309L type electrode (weldment B) and a 310 typeelectrode (weldment C). The welding parameters aregiven in Tables 2-4. These tables also show the weldpreparation, welding sequence, a macrograph and the

chemical analysis of deposited weld metal.

Weldment AWeld preparation Welding sequence

Weldment B

Weld preparation Welding sequence

Weldment C

Weld preparation Welding sequence

Table 2-4: Welding parameters

Weldment A Weldment B Weldment C

Base material Avesta 253 MA Avesta 253 MA Avesta 253 MA

Plate thickness 12 mm 12 mm 12 mm

Filler material Coated electrode - 253 MA Coated electrode E309-16 Coated electrode E310-16

( = 3.25 and 4 mm)Preheat No No No

Heat input 1.3 kJ/mm 1.3 kJ/mm 1.2 kJ/mm

Chemical analysis (weight percent)

C Mn Si S P Cr Ni Ce N

Base material 0.09 0.56 1.52 0.001 0.024 20.8 11.2 0.05 0.179

Weldment A

Deposited weld metal 0.039 0.77 1.68 0.004 0.015 21.0 10.2 ND ND

Weldment B

Deposited weld metal 3.25 0.017 1.14 0.80 0.014 0.015 24.3 12.7 ND ND

4.00 0.021 1.16 0.72 0.015 0.014 23.5 12.8 ND ND

Weldment C

Deposited weld metal 0.09 2.60 0.73 0.007 0.011 25.0 19.8 ND ND


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Tensile test resultsTensile tests are performed at a range of temperaturesbetween room temperature and 1000C on both basemetal and on welded specimens taken perpendicularto the welding direction (cross weld specimens). Thebase metal tensile tests are performed on an electro-mechanical test machine with furnace heating. Theweldments are tested both in a servo-hydraulic tensilemachine in which the specimen is resistance heated(weldment A and B) and in an electromechanicalmachine with furnace heating (weldment C). Resis-tance heating of the specimen has the disadvantageof creating a thermal gradient along the specimen'saxis, but shows the advantage of probing the actualweld metal properties instead of those of the weldmentas a whole. In the case of the furnace heated weld-ments C, however, the failure location was also situatedat the weld metal occupying the minimum cross sec-tion.

The test results are listed in Tables 5 and 6 and shownin Figures 5 and 6 (page 6).

Significant differences in strength properties are notedbetween the base material and the weld metals. Thematching weld metal A and weldment C have highertensile strengths than weld metal B, but all weld metalsshow ultimate tensile strengths inferior to that of thebase metal. In contrast, the 0.2% yield strength of the310 type weldment C is slightly higher than that of thebase metal.

The base metal shows a pronounced ductility dip in thetemperature range 650-700C, while all weld metalsexhibit much lower high temperature ductility: for thereduction in area this applies at temperatures above700C, while for the elongation at all test temperatures.

Table 6: Tensile test results (weld metal)Strain rate: 1 10

-2mm/sec

Test 0.2% Ultimate Elonga- Reduc-

temp. Yield tensile tion tion

strength strength in area

(C) (N/mm2) (N/mm

2) (%) (%)

Weldment A:

400 - 550 30.8 48.2

500 - 500 24.8 46.7600 - 437 16.8 39.2

700 - 342 10.0 32.8

800 - 235 4.8 17.2

900 - 125 3.6 11.6

1000 - 70 10.8 12.8

Weldment B:

600 - 337 12.0 46.7

700 - 270 6.0 40.7

800 - 165 4.0 26.0

900 - 90 - 19.0

1000 - 58 5.6 17.8

Weldment C:

20 393 630 20.5 30.6500 191 478 22.5 26.2

600 179 435 18.5 32.7

700 177 314 9.0 19.5

800 163 192 6.3 17.1

900 115 129 7.5 19.5

1000 51 58 7.5 10.0

800Table 5: Tensile test results (base material)Strain rate: 1 10

-2mm/sec

Test 0.2% Ultimate Elonga- Reduc-

temp. Yield tensile tion tion

strength strength in area

(C) (N/mm2) (N/mm

2) (%) (%)

20 330 715 60.5 73.7

100 280 653 52.2 73.7

200 235 603 47.0 71.0

300 201 597 47.5 68.1

400 196 593 49.3 68.1

500 172 555 45.0 62.1600 168 498 43.0 58.9

650 153 465 43.8 34.4

700 165 365 31.8 40.8

700 143 378 46.7 38.4

800 151 229 78.8 52.0

800 400 242 78.8 52.0

900 113 118 51.6 90.5

900 98 161 84.0 68.0

1000 67 61 62.0 93.4Figure 5Tensile test results (Strength)


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Figure 6Tensile test results (Ductility)

Low strain rate tensile

testing (LSRTT)Experimental procedure:

For each configuration a series of specimens is testedwith strain rates ranging from 10

-5to 10

-8mm/mm sec.

Therefore, cylindrical specimens are taken out of thebase material or over the weld to be tested andmachined according to Figure 7. In order to avoidbending, the specimens are provided with very accu-rately machined endpieces for centering.

During the tests, the specimens are heated either byresistance heating or by furnace heating in a three zonesplit furnace. Temperature is controlled by a thermo-couple, spotwelded in the centre of the specimen. Thefirst heating method allows very fast heating of thespecimen, but it has the disadvantage of a non-uniformtemperature distribution along the gauge length of the

specimen.

The elongation of the specimen is measured with a hightemperature extensometer, provided with two quartzrods which are mounted on the specimen over a 25 mmgauge length. When resistance heating is applied, thedeformation, calculated from the actual gauge length,has to be corrected for the thermal gradient along thespecimen's axis.

Testing is performed either in a servo-hydraulic or in anelectromechanical closed loop tensile testing machine;the tests are run under displacement and strain controlrespectively. The control signal is produced by a preci-sion low frequency ramp generator. A load-deformation

diagram is continuously registered; the tests arestopped when the load reaches a steady state value(Figure 8). It is normal practice that tests are carriedout over four decades: 10

-5, 10

-6, 10

-7and 10

-8

mm/mm sec; nominal testing times range from 2 hoursto about 3 days.

Figure 7LSRTT test specimen (dimension in mm)

Figure 8Low strain rate tensile test curve


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Interpretation of low strain rate tensile testresults:

In low strain rate tensile tests, the load taken up by thespecimen reaches a saturation level that depends onthe imposed strain rate (fig. 8). This steady state condi-tion is a manifestation of a dynamic equilibrium situ-ation, comparable to the secondary or steady statecreep stage in a creep test under a constant applied

stress (fig. 9). The saturation stress s in LSRTT can beconsidered as the resistance of the material againstslow straining, and thus against creep deformation.Although the combinations of stress and strain rate increep and low strain rate tests are not equivalent (2),the results of LSRTT can be used confidently for creep-life prediction purposes (4). The main advantage ofLSRTT is that the testing can be performed in a veryshort time period.

Low strain rate tensile tests at 800C have been per-formed on 253 MA base material and cross weld speci-mens from weldments A and B. For comparison pur-poses, tests were also performed on alloy 800H basematerial. The results are listed in table 7 and plotted

in figure 10. From this figure it can be observed thatthe saturation or steady state stress of the alloy 800Hbase material is lower than that of the 253 MA basematerial, while the saturation values for both weldmentsare still higher. Consequently, the following order increep strength is expected: weldment A >weldmentB > base material 253 MA > base material alloy 800H.

Figure 9Creep curve

Table 7: Low strain rate tensile test results at 800C253 MA and Alloy 800H

(s-1) s (N/mm

2)

Base metal 253 MA 10-5

150.2

10-6

109.0

10-7

68.3

10-8

45.8

Weldment A 10-5

165.0

10-6

128.1

10-7

87.7

10-8

63.2

Weldment B 10-5

140.6

10-6

104.0

10-7

88.7

10-8

48.9

Base metal Alloy 800H 10-5

97.0

10-6

78.3

10-7

51.2

10-8

37.9

Metallurgical stabilityIt is well known that austenitic stainless steels afterlong-time exposure at high temperature (> 600C)can give rise to sigma-phase formation. In the case of253 MA, the formation of Pi-phase (Cr12Si12Fe4Ni2N4)is also reported (1). These hard and brittle phasesusually adversely affect the mechanical properties,particularly the impact strength at ambient tempera-ture; the toughness at higher temperatures may remainsatisfactory.

Welded test plates (Weldment C) were artifically agedup to 1000 hrs at 500, 600, 700, 800 and 900C, a

second series was aged up to 10,000 h at 550C only.Subsequently, the mechanical properties were deter-mined in conventional tensile testing at 20C and at theageing temperature; the room temperature toughnesswas evaluated by means of Charpy-V notch impacttesting. Figure 10

Low strain rate tensile test results at 800C


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Figures 11 and 12 respectively show the room tempera-ture strength and ductility after 1000 hrs ageing at theabove mentioned temperatures. Both base metal andweld metal strength remain approximately constantafter ageing at all temperatures. Although the weldmetal has initially a lower ductility than the base metal,comparable values are found when tested at room tem-perature after ageing at temperatures above 700C.For the base metal, there is a significant decrease of

elongation and reduction in area after 1000 hrs ageingat 700C; this loss of room temperature toughness,which is most pronounced at 800C, is also observedby impact testing (Figure 13).

Figure 14 shows the elevated temperature strength andductility after 1000 hrs ageing. Strength values for bothbase metal and weld metal, when tested at the ageingtemperature, hardly deviate from the results of theunaged base and weld metal (Figure 5), obtained pre-viously.

With regard to ductility, the following trends are found:

a. after 1000 hrs ageing at temperatures above 600C,an increase in the elevated temperature ductility of

the base material is observed.b. for weld metal, however, the elevated temperature

ductility decreases to 10-15% at 700C and re-mains rather poor up to 900C.

The influence of different ageing treatments on theroom temperature impact strength of both base metaland weld metal C is shown in Figure 15. Even long-timeageing temperatures as low as at 550C result in aserious loss of toughness of the base metal; fromaround 200 J in the as-received condition down toabout 40 J after 10,000 hrs.

Microstructuraland fractographicexaminations

(253 MA Base Material)The techniques employed in this study are:

a. microstructural examination by lightmicroscopy(LM) and transmission electron microscopy (TEM).The specimens for optical investigations weremechanically polished using conventional polishingtechniques. The general microstructure was re-vealed by the mild Kalling reagent, sigma-phase waselectrolytically revealed with 1N KOH. For TEMinvestigations both extraction replicas and thin-foilswere prepared.

b. scanning electron microscopy (SEM) and TEM of thefracture surfaces from Charpy-V impact specimens.For the latter carbon extraction replicas were pre-pared by using the direct stage carbon extractiontechnique.

c. determination of the composition of the phases byelectron diffraction in a 120 kV TEM and by EDAX-analysis in SEM and TEM.

Figure 11Influence of 1000h ageing on the room temperature strength

Figure 12Influence of 1000h ageing on room temperature ductility

Figure 13Influence of 1000h ageing on the Charpy V notch impact

strength at room temperature


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Figure 14Influence of 1000h ageing on the elevated temperaturestrength and ductility of 253 MA base material (Test tempera-ture = ageing temperature)

Figure 15Influence of different ageing treatments on the room tempera-ture impact strength


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Figure 16Influence of artificial ageing on the microstructure of 253 MA- x 400

Figure 17SEM-pictures - fracture surface of Charpy-V test specimenartificially aged for 1000h at 700C

As received 1000h - 700C

1000h - 500C 1000h - 800C

1000h - 600C 1000h - 900C

A: x 250 B: x 550

C: x 400 D: x 2000


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The microstructural evolution after ageing for 1000 hrs.at different temperatures ranging from 500 to 900C isshown in Figure 16. Precipitation of carbides is clearlyvisible after ageing at 600C, the M23C6-carbides beingmainly located at grain boundaries and along the crys-tallographic planes of the matrix.

After ageing at 700C the grain and twin boundaries areheavily decorated with coarse precipitates while fineprecipitates are also found throughout the matrix.These precipitates on grain and twin boundaries andin the matrix were identified as Cr23C6 and sigma-phase. Nitrides could not be detected. SEM-investiga-tions of the Charpy-V notch fracture surface of materialaged at 700C show a mainly intergranular fracture with

randomly distributed flat areas (Figure 17). These flatareas were recognized as the fracture-path through theaustenite/sigma-phase interphase of the twin bound-aries, because of their favourable orientation in relationto the impact direction.

After 1000 hrs ageing at 800C a coalescence of pre-cipitates was found. On the grain and twin boundaries,Cr23C6-carbides, some Pi-phase and an almost con-tinuous network of sigma-phase could be observed andidentified (Figures 18 and 19). The Pi-phase could bediscriminated from sigma-phase by EDAX-analysis froma higher Si-content of the former. The fracture surfaceshows an almost intergranular-ductile aspect withlocally some flat areas (Figure 20).

Figure 18Optical micrograph - sigma-phase in 253 MAafter artificial ageing for 1000h at 800C etchedin INKOH

Figure 19TEM-picture - carbon extraction replica from Charpy-V frac-ture surface (253 MA artificially aged for 1000h at 800C)Fracture through sigma-phase x 10,500

Figure 20SEM-picture - fracture surface of Charpy-V test specimenartificial ly aged for 1000h at 800C

A: x 570 B: x 1000

C: x 1600 D: x 6500


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Figure 21SEM-picture - fracture surface of Charpy-V test specimenartificially aged for 1000h at 900C

Figure 22Influence of artificial ageing at 550C on the microstructure- x 500

A: x 150 B: x 600

C: x 775 D: x 3100

A: 1000h at 550C B: 10,000h at 550C


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After ageing for 1000 hrs at 900C, coarse precipitatesare found at the grain and twin boundaries (coherentand incoherent). The discontinuous network of pre-cipitates on the grain boundaries could be clearly identi-fied as sigma-phase by electron diffraction analysis.

SEM-investigations on the fracture surface show a com-bination of more ductile features (dimples) and inter-granular decohesions (Figure 21).

Long time exposure (10,000 hrs) at 550C shows apronounced grain boundary and a slight twin boundaryprecipitation (Figure 22B). The Charpy fracture sur-faces show mixed intergranular and ductile features(Figure 23). TEM-investigations on carbon-replicasshow the presence of only M23C6-carbides while noembrittling phases, i.e. sigma or Pi, could be revealedin this specimen (Figure 24).

DiscussionThe fully austenitic stainless steels are generally sen-sitive to hot-cracking; however, the hot ductility testsshowed that 253 MA material in the as-received condi-tion is not sensitive to the formation of heat-affected-zone hotcracks. Subsequent test welds with threedifferent types of coated electrodes confirmed thisresult; indeed no welding problems such as (hot-)crack-ing were experienced.

253 MA base metal, in the as-received condition, hasa high strength at elevated temperature; short-timeelevated temperature tensile testing shows even betterresults than given in the technical documentation. Theductility at elevated temperatures remains on a highlevel although in the temperature range 600-750C aductility dip was found. However, both reduction in areaand elongation are still acceptable (>30%) and su-perior to those of all investigated weld metals.

A: x 570

Figure 23SEM-picture - fracture surface of Charpy-V test specimenartificially aged for 10,000h at 550C (253 MA base metal)-x 570

Figure 24TEM-picture - carbon extraction replica from Charpy-V frac-ture surface (253 MA artif icially aged for 10,000h at 550C)-x 3900

B: x 570


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Limitations in application will be imposed by the lowductility at elevated temperatures of the weld metal, ifa 309, 310 or a matching type of filler metal is used.

After ageing for 1000 hrs at several temperatures, theroom temperature and high temperature strength ofbase metal and weld metal are almost unaffected.However, the room temperature ductility of base andweld metal has strongly decreased after ageing attemperatures above 550C and most importantly at800C; where low figures for ductility (15%) and ab-sorbed energy (10 Joule) are obtained.

The room temperature embrittlement after ageing attemperatures of 700, 800 and 900C can be attributedto the formation of sigma phase which leads to mainlyintergranular fractures. Also Cr23C6-carbides wereformed and found along the grain boundaries and thosewill also contribute to the embrittlement. This isevidenced by the fact that a serious loss of impactstrength due to M23C6 precipitation only, is found afterextended ageing at 550C.

Other embrittling phases such as Pi-phase are foundin very small amounts after ageing at 800C only.

From these observations it is possible to summarizethe effect of ageing on the room temperature impactstrength in a time-temperature diagram (fig. 25). In thisdiagram an "iso-embrittlement" corresponding to aroom temperature impact value of 30-50 J can bedrawn. It should be noted that this "embrittlement" isfound after ageing times as low as 100 hours and notonly after ageing for over 2000 hours (5). An interestingfeature is the stability of impact strength (about 50-60Joule) after ageing between 10 and 1000 h at 900C.Although sigma-phase is known to be present, the roomtemperature impact strength remains relatively high;

coalescence of sigma-phase and a discontinuous net-work was found. These features seem to be very bene-

ficial for the room temperature toughness. Because ofthis, a further decrease in room temperature impactstrength and toughness is not expected, even afterlonger exposure times.

The Low Strain Rate Tensile Testing-method has provedto be a useful tool in selecting the appropriate fillermetal (and welding procedure) in order to obtain thedesired weldment creep strength. The results clearlyindicate that a weldment with the (253 MA) matchingelectrode has the highest creep strength and is evensuperior to base metal. In the case of this matchingelectrode, however, no check was made on the metal-lurgical stability as no testing of weld metal after arti-

ficial ageing has been performed. Nevertheless, onecan expect that comparable results as with 253 MAbase metal will be obtained after ageing and mech-anical testing.

A weldment with the 309 type of electrode howevershows a lower internal stress level than that of the weld-ment with the matching 253 MA electrode and has thesupplementary disadvantage that at lower strainratescreep strength will fall below the base metal creepstrength. This is indicated by the 309-weldment curvewhich converges with the base metal curve at thoselower strainrates (Figure 10). This is most probablycaused by the low carbon content of this weld metal( 0.021% only).

Although LSRTT is not performed on the 310-weldment,it can be expected that due to the higher carbon con-tent a higher creep strength than that of the base metalwill be obtained. On the other hand, a serious loss ofductility and impact strength was found for this weldmetal; although not examined, it is anticipated thatsigma-phase formation is the main cause for this em-brittlement.

Therefore, weld metal selection for high temperatureservice and applications shall be based on compro-mises between creep strength and metallurgical sta-bility, especially embrittlement in the lower tempera-ture range.

Remarkable is the lower creep strength (evidenced bylow strain rate tensile testing) of the Alloy 800Hcompared to the 253 MA material. The only reasontherefore is that the low grain size of the tested Alloy800H (ASTM 7-8) which is due to solution annealingat too low a temperature (1150C). This highlights theneed for a correct solution annealing temperature(1175C) to obtain the appropriate grain size (ASTM4-5) of the Alloy 800H for high temperature applica-tions (6-7).

Figure 25Time-temperature diagram (Influence of ageing) encirclednumbers are impact values (J) at room temperature

Conclusions1. The fully austenitic stainless steel UNS S30815

(253 MA) is not sensitive to heat affected zone hot-cracking.

2. In high temperature applications i.e. 500 to 900C.room temperature embrittlement due to sigma-phase formation should be taken into account.

3. At 800C, the creep strength of 253 MA is superior

to AISI type 304H and even better than that of thetested Alloy 800H (solution annealed at 1150C).

4. Selection of the appropriate weld metal must bebased on a compromise between creep strengthand metallurgical stability. The matching 253 MAelectrode shows the most promising results.


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AcknowledgmentsThe authors acknowledge the financial support of B.I.L,DOW Chemical Nederland B.V., I.W.O.N.L, N.F.W.O.and W.T.C.M.

Thanks are due to the staff of the Laboratorium Soetevoor Weerstand van Materialen en Lastechniek, Uni-versity of Gent, the Materials Investigation Depart-

ment of DOW Chemical Nederland and the Laboratoryfor Electronmicroscopy (dr. ir. K. Ostyn), where theexperimental work has been carried out. Also ourthanks to S.A. SOUDOMETAL - Brussels for weldingpart of the test panels. Guidance and suggestions byIr. F.J. Vaes and Prof. Ir. A. Vinckier are gratefully ac-knowledged.

Avesta is very much obliged to the authors for their kindpermission to publish this report and to the Research

Center of the Belgian Welding Institute where theresearch work was supervised.

References1 Technical documentation Avesta, Sandvik.

2 Steen M., Creep life assessment by low strain ratetensile testing. Int. J. Pres. Vess. & Piping, 14, 1983.

3 Kreischer C.H., A critical analysis of the weld heat-affected-zone hot ductility test. Welding ResearchSupplement, Feb. 1963.

4 Provost W., Steen M., Dhooge A., Bepaling van dekruipeigenschappen van hoge temperatuursmate-rialen d.m.v. korte duurproeven (Creep life assess-ment by low strain rate tensile testing). Lastechniek,November-December 1984.

5 Andersson T., Microstructure and properties of thetwo high temperature steels SANDVIK 253 MA andW.-Nr. 1. 4828. SANDVIK Technical Publication, (Jan.1979).

6 Hoek W., The notch sensitivity of Alloy 800H undercreep conditions. To be published.

7 Degischer H. P., Aigner M., Lahodny H., Spirader K.,

Qualification of stationary creep of the carbide pre-cipitating Alloy 800H. Conference on High Tempera-ture Alloys. Oct. 15-17, 1985, Petten, Nederland.

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