Intergranular Attack in Welded Stress-Corrosion Resistant Stainless Steels

Satisfactory stress corrosion resistance is obtained using arc welding and, depending on the base metal,

Type 376/. or 309L stainless steel filler metals

BY J. HONEYCOMBE AND T. C. GOOCH

Introduction

Although austenitic stainless steels have excellent corrosion resistance in a wide variety of media, they exhibit a high susceptibility to stress corrosion cracking (SCC) in hot aqueous envi­ronments containing chloride or hy-droxyl ions. Industrial interest has, therefore, been increasing in ferritic and duplex-structured stainless steels (containing 17-27% chromium with 0-5%Ni) which are resistant to stress corrosion cracking. Despite their resistance to SCC, the conventional ferritic stainless steels containing about 17%Cr possess several disadvan­tages; in particular, their welded joints may show excessive grain growth and embritt lement, and sensitivity to inter­granular corrosion.'

Alloys have been recently devel­oped in an attempt to overcome these problems. In view of the wide range of potential applications of such alloys, an investigation was initiated at The Welding Institute to study their weld­ing behavior so that practical recom­mendations regarding opt imum pro­cedures can be formulated. The pres­ent paper considers the effects of welding in causing susceptibility to intercrystalline attack of representa­tive materials.

Materials

The alloys studied were intended to cover the different microstructural

Paper presented at the 58th AWS Annual Meeting held in Philadelphia, Pennsylvania, during April 25-29, 1977.

j. HONEYCOMBE is Senior Research Met­allurgist and T. C. COOCH is Principal Research Metallurgist, The Welding Insti­tute, Abington, Cambridge, England.

approaches to avoiding chloride-induced SCC. The materials were a conventional ferritic grade (AISI 430) to afford a base-line for comparison, a ferrite-martensite alloy, two ferrite-austenite steels, and a fully ferritic steel having extra low interstitial (ELI) element content.* The particular ma­terials employed and the available thicknesses are given in Table 1.

Test Weld Preparation

The processes employed were shielded metal arc (SMA), gas tungsten arc (GTA), gas metal arc (GMA), submerged arc (SA) and electron beam (EB). Apart from the autogenous GTA and EB welds, three types of commer­cially available filler metals were selected for study—a matching or near-matching alloy; an austenitic stainless steel (either 18/10/3Mo or 23/12/Mo, subsequently referred to as Types 316L or 309L-Mo respectively) as might be employed in practice;- and a nickel alloy, as in principle representing a weld metal of high resistance to chlo­ride-induced stress corrosion. The nominal compositions of the filler metals used are shown in Tables 2 and 3, w i th available deposit analyses.

Butt welds were made in a restrain­ing jig as indicated in Tables 4 and 5, wi th interpass temperature being held at 150 C (302 F) maximum. No weld was given a postweld heat treatment. Transverse sections of all welds were

*Material was tested as 2.5 mm (0.1 in.) thick sheet and 13 mm (0.5 in.) plate. It is understood from the material manufacturer that subsequent to carrying out the present tests on the thinner material, the composi­tion has been modified, although no details are available. The thicker material was supplied and tested after the reported composition change.

examined metallographically.

SMA Welds

All the manual metal arc welds employed a 75 deg included angle, single V-groove preparation with a 1.5 mm (0.060 in.) root face and no root gap, and were made wi th an autoge­nous GTA root and argon back purging for better control of penetration.

The welds in the ferritic-austenitic alloys were deposited using nomi­nally high and low energy inputs.t With the single phase 26%Cr-1%Mo, because of its tendency to grain growth, low energy inputs only were used.

GTA Welds

Joint preparation and root run were as used with the SMA welds. For the 26%Cr-1%Mo steel, the argon gas f low through the torch was kept high at approximately 25-30 l /min wi th back purging maintained at 6 l /min; 1 for the other steels, the torch f low rate was 10 l /min.

GMA Welds

The preparation used was a 70 deg single V-groove with a 3 mm (0.12 in.) root face and no root gap. Following welding from one side, the weld was back gouged and a sealing run inserted on the reverse side.

The shielding gas for the two duplex alloys was argon-2%0L. or argon depending on the filler, whi le for the 26%Cr-1%Mo it was either pure argon or argon-1%0,. The gas f low was kept high with argon back purging in all

jEncrgy input defined as: (Volts X Amps X 60/speed (mm/min) X 1000) x kj/mm

W E L D I N G RESEARCH SUPPLEMENT I 339-s

Table 1-Analyses of Steels Studied, Wt-%

Sample

Steel 1: 17%Cr Steel 2: 16%Cr-2.5%Ni-

1%Mo Steel 3: 26%Cr-1%Mo

Steel 4: 26%Cr-5%Ni-1.5%Mo

Steel 5: 18%Cr-5%Ni-2.5%Mo

Material thickness, mm

2.5

2 2.6

13

10

13

C

0.083

0.035 0.007 0.003

0.026

0.025

Mn Si S P Ni Cr Mo

0.33 0.34 0.009 0.029 0.23 16.44 0.08

Nb Ti Cu O N

0.005 <0.01 0.08 0.006 0.034

0.66 0.47 0.006 0.009 2.72 15.91 1.22 0.5 <0.01 0.03 0.006 0.028 0.007 <0.01 0.23 0.009 0.01 0.003 <0.05 0.35 0.009 0.01

0.74 0.63 0.006 0.02

1.66 1.69 0.010 0.02

0.12 25.8 0.15 26.2

0.99 0.005 <0.01 <0.01 0.002 0.012 0.99 0.06 <0.05 <0.05 0.002 0.012

5.15 26.25 1.3

5.2 18.7

0.05 <0.05 <0.05 0.005 0.07 <0.05

2.78 <0.05 <0.05 0.003 0.07

Table 2-Nominal Analyses of Filler Metals, Wt-%

Filler metal C Cr Mo

Matching steel 3 bare <0.005 27.5 0.99 electrode

Matching steel 4 0.10 26.0 1.50 covered electrode 0.05 25.5 2.0

Type 316L covered 0.03 19.0 3.0 electrode max

Type 309L-Mo covered 0.03 23 2.5 electrode

Ni-base covered 0.1 13- 0.5-electrode max 17 2.5

Near matching steel 4 0.06 25 2 bare electrode

Type 316L bare 0.025 18.5 2.7 electrode

Ni-base bare electrode 0.030 20 max

Ti

-

—

-

-

-

-

-

0.3

Si

0.22

0.50 0.5 0.25

0.8

0.75 max 1.0

0.85

0.2

Mn

0.02

1.50 1.4 1.3

2.5

2

1.0

1.35

3.0

Ni

0.06

5.0 6.5

12.5

12

Bal

5

11.5

72

Nb S P Fe

- 0.014 0.010 Bal

-

-

-

1.0-3.0

-

-

2.5

0.015

0.02

0.03

0.02 max

-

0.015

0.010 max

0.015

0.03

0.04

—

-

0.015

0.010 max

Bal

Bal

Bal

6.12

Bal

Bal

1.0 max

Cu N

0.02 <0.012

cases. W e l d i n g was m e c h a n i z e d , a n d t h e t o p o f each run was g r o u n d t o r emove any o x i d e c o n t a m i n a t i o n .

SA Welds

Incoflux 6, dried for 2 h at 200 C (392 F) was used for all welds. Joint prepa­ration was as for the GMA welds. Welding on the first side was com­pleted, the root run removed and a sealing run deposited.

EB Welds

The EB welds were made wi th a square butt preparation and no filler metal.

Corrosion Testing

General Comments

A number of different test methods have been developed to assess suscep­tibil ity to intercrystalline attack in stainless steels. When selecting evalu­ation techniques, it is necessary that the data obtained should be relevant to practice and, as far as possible, should assess behavior under a range of conditions. In the present work, the corrosion resistance of the test welds was examined in two ways—namely, by using the acid copper sulphate (Strauss) test and by using a potent io-static technique.

The Strauss test was selected in

preference to other comparable im­mersion tests, such as the acid ferric sulphate procedure, on the basis of its generally satisfactory history of prac­tical application for other than highly oxidizing media. Another reason was its common use in the UK for quality control of stainless steel products.

The potentiostatic technique was employed because, in principle, it gives information relevant to a range of environmental condit ions and has been successfully used to define conditions causing intercrystalline at­tack in austenitic materials.'

Strauss Testing

Strauss tests were conducted on all welds according to BS 1449: Part 4: 1967. Three coupons were removed transversely to the welds, as in Fig. 1, and the sides abraded to a 600 emery finish.

One coupon was used as a control sample to allow for any difficulty in interpreting results of the necessary bend tests due to sample brittleness. The control specimen was bent 90 deg around a former of radius three times the plate thickness.

The specimens were considered to have failed the bend test if cracks appeared in the weld or heat-affected zone (HAZ). Any cracks appearing at the cut edges of the specimen were ignored.

The remaining two coupons were suspended for 72 h in boi l ing acidified copper sulphate; on removal one was bent as for the control specimen. From the second, a section was taken trans­verse to the weld, mounted and polished. The depth of penetration of any grain boundary attack was mea­sured at X400 magnification."'

In addit ion, selected welds were exposed to the "mod i f i ed " acid cop­per sulphate test electrolyte w i th copper turnings in the test flask." Susceptibility to intergranular attack was assessed by depth of penetration measurements.

Potentiostatic Studies

In the first instance, potentiostat scans were conducted on all base metals under investigation. Specimens with 1 cm2 surface area were mounted and polished to a 1,um finish. The basic electrolyte used was a 20% H,SO, containing O.lg/I NH..CNS and the applied voltage was changed at 0.3 V /h .

From these scans, the approximate potentials at which the steels became passive were identif ied. Particular ref­erence was made to the potential range corresponding to moderately oxidizing conditions, since corrosion in service wi l l most frequently involve a passive/active transition in this potential range.7

340-s I NOVEMBER 1977

Table 3—Deposit Analys

Consumable

Matching steel 3 bare electrode

Matching steel 4 covered electrode

Type 316L covered electrode

Type 309L-Mo covered electrode

Nickel-base covered electrode

Near matching steel 4 bare electrode

Type 316L bare electrode

Nickel-base bare electrode

:es, Wt-%'

Base metal

Steel 3 Steel 3 Steel 3

Steel 4

Steel 4

Steel 3 Steel 5

Steel 3 Steel 4 Steel 5

Steel 4 Steel 4 Steel 4

Steel 4 Steel 5 Steel 4 Steel 5

Steel 3 Steel 4 Steel 4 Steel 5

•'

Process

GTA G M A " "

SA

SMA

SMA

SMA SMA

SMA SMA SMA

GTA GMA

SA

GMA GMA

SA SA

SA GTA

SA SA

C

0.006 < 0.026

0.04

0.08

0.05

0.02 0.02

0.04 0.04 0.04

0.05 0.05 0.06

0.02 0.02 0.04 0.04

0.03 0.03 0.04 0.04

Cr

28.2 27.1 26.1

26.9

21.0

24.4 22.9

13.8 16.3 15.6

26.5 24.7 26.2

19.7 18.8 20.5 18.2

22.0 20.9 22.1 21.3

M o

0.90 0.98 1.01

1.46

2.62

2.36 2.7

1.13 0.90 1.12

2.31 2.2 2.14

2.59 2.81 2.38 2.75

0.29 0.09 0.50 0.60

Ti

< 0 . 0 5 < 0 . 0 5 < 0 . 0 5

< 0.05

< 0.01

0.01 0.05

< 0.05 < 0 . 0 5 < 0 . 0 5

< 0.01 < 0 . 0 1 < 0.05

< 0.01 < 0.01

0.01 0.02

0.14 0.36 0.09 0.17

Al

N D ND

< 0.05

< 0.05

< 0.05

0.05 0.05

< 0.05 < 0.05 < 0.05

< 0.05 < 0.05 < 0.05

< 0.05 < 0.05 < 0.05 < 0.05

0.08 < 0.05

0.10 0.07

Si

0.20 0.22 0.31

0.98

0.22

0.63 0.89

0.27 0.25 0.39

0.39 0.38 0.48

0.77 0.83 0.78 1.01

0.56 0.41 0.67 0.84

Mn

< 0.05 < 0.05

0.05

0.76

1.28

1.36 1.68

2.23 1.57 1.96

0.86 0.83 0.81

1.26 1.40 1.38 1.59

2.43 2.94 2.28 2.92

Ni

0.07 0.14 0.13

5.14

1 1.0

10.0 11.1

70.4 63.1 63.2

5.1 4.7 5.35

10.8 11.0 10.7 11.0

55.0 66.9

Bal 59.2

Nb

0.06 0.07

< 0 . 0 5

< 0.05

< 0.01

< 0.05 < 0.05

2.28 1.82 2.09

< 0.01 < 0.01 < 0.01

< 0.01 < 0 . 0 1 < 0.01 < 0.01

1.80 2.36 1.72 1.95

S

0.009 0.006 0.007

0.012

0.015

0.013 0.013

ND ND ND

0.016 0.017 0.015

0.012 0.012 o.ori 0.012

ND ND

0.012 ND

P

0.011 0.016 0.016

0.030

0.016

0.026 0.020

ND ND ND

0.026 0.025 0.026

0.014 0.014 0.016 0.017

ND ND

0.020 ND

Fe

Bal Bal Bal

Bal

Bal

Bal Bal

9.6 15.4 15.4

Bal Bal Bal

Bal Bal Bal Bal

17.6 5.9

26.1 14.7

Cu

0.05 < 0 . 0 5 < 0 . 0 5

< 0.05

0.03

0.05 0.06

0.13 0.09 0.11

1.4 1.32 1.04

0.02 0.03 0.03 0.05

0.05 0.06 0.1 0.07

' "ND = nol dele ""Ar shielding.

The critical potentials below which corrosion occurred were more precise­ly defined for the various weld areas by a controlled potential etching procedure.7 The polished specimens were exposed to the electrolyte at a potential slightly more positive than the passivation potential indicated by the scans, and examined at X500 magnification for any sign of corro­sion. If none had taken place, the potential was lowered by 0.01 V inter­vals, the specimen being repolished before each change, until attack was first observed.

The nature of the attack was identi­fied and the potential below which preferential intercrystalline weld metal or HAZ attack occurred was noted. Preliminary assessment was based on a 20 min exposure period; critical poten­tials were then more closely defined using 16 hr and 64 hr periods.

Additional studies were carried out in the basic electrolyte wi th 3% NaCI additions. This was for two reasons: First, trials in 20% H,SO, + 0.1 g/l NH4CNS indicated the passive fi lm formed on the experimental steels to be kinetically stable, and the effect of chloride ions in promoting passive film breakdown was examined. Second, intergranular attack in service has been associated with chloride media;s it

Table 4—Conditions for GTA Ri

Steel no.

1 2 3

'"25.4 mm = I inch

Volts, V

11 10 12

ins in Thin Sheet Samples

Currents, A

120 100 100

Travel speed,

m m / m i n

190 190 125

Input energy,

k j / m m " "

0.4 0.3 0.5

was, therefore, desirable to examine the role of chloride ions in inducing preferential grain boundary corrosion in the materials studied.

Results

Metallographic Observations

Steel 1: AISI 430 (17%Cr). The base metal was fully ferritic containing bands of chromium carbide particles. After autogenous GTA welding, both the weld metal and the HAZ exhibited duplex structures wi th martensite along the grain boundaries and on certain planes throughout the grains—Fig. 2.

The grain size of the weld metal was about ASTM 1. Close to the fusion boundary where transformation to austenite had occurred during weld­ing, the HAZ exhibited a smaller grain

size closer to that of the base metal. In the low temperature HAZ, the struc­ture reverted to single-phase ferrite wi th a grain size (ASTM 5) only slightly larger than the base metal (ASTM 6).

5tee/ 2: Ferrite-Martensite (16%Cr/ 2.7%Ni/1%Mo/0.5%Nb). The base metal showed a fine, banded ferrite-martensite structure, containing Nb(C,N) particles—Fig. 3. In the weld metal and high temperature HAZ of autogenous GTA welds, transforma­tion occurred to a fully ferritic struc­ture of ASTM 3-4 grain size, wi th Nb(C,N) precipitation at grain bound­aries and throughout the grains. Nio­bium carbide eutectics were apparent in the HAZ.

Steel 3: Ferritic ELI (26%Cr/1%Mo). The base metal was fully ferritic—Fig. 4. In both thicknesses examined, the grain size was variable, giving a banded appearance; in the 13 mm (0.5

W E L D I N G RESEARCH SUPPLEMENT I 341-s

Table 5-

Weld no.

1

2 3 4

5

6 7

8

9 11 12

2

3 4 5 6 7

8

9 10

11 12 13

14 15

1 2 3 4 5 6 7

8 9

10 11 12 13

""One l/m ""25.4 mm

-Welding I

Process

SMA

SMA SMA GMA

GMA

GMA GTA

GTA

GTA SA EB

SMA

SMA SMA SMA SMA GMA

GMA

GMA SA

SA SA GTA

GTA EB

SMA SMA SMA SMA SMA SMA GMA

GMA SA SA GTA GTA EB

Parameters for Plate Samples

Filler metal and diameter,

mm

Type 309L-MO 3.25

Type 316L 3.25 Ni-base 3.25 Matching composi­

t ion 1.6 Matching composi­

t ion 1.6 Ni-base 1.2 Matching composi­

t ion 2.3 Matching composi­

tion 2.3 Ni-base 1.6 Ni-base 1.6

-

Matching composi­t ion 3.25

Type 316L 4 Type 316L 3.25 Ni-base 4 Ni-base 3.25 Near matching

composit ion Type 316L 1.2

Ni-base 1.2 Near matching

composition 1.6 Type 316L 1.2 Ni-base 1.6 Near matching

composit ion 1.6 Ni-base 1.6

-

Type 309L-MO 4 Type 309L-MO 3.25 Type 316L 4 Type 316L 3.25 Ni-base 4 Ni-base 3.25 Type 316L 1.2

Ni-base 1.2 Type 316L 1.2 Ni-base 1.6 Type 316L 1.2 Ni-base 1.6

-in = 2.1 cu fl/hr. = 1 inch.

"'Excluding autogenous GTA root runs.

Shielding gas" '

or flux

Steel 3,

---Argon 30 l /min

Argon + 1 % 0 , 30 l /min

Argon 25 l /min Argon 20 l /min

Argon 25 l /min

Argon 30 l /min Inco flux 6

-

Current, A

26%Cr/1%Mo 90

85 80

250

250

220 140

220

145 270

45 x 1 0 "

Steel 4, 26%Cr/5%Ni/1.5%Mo

-

----Argon + 2%02

25 l /min Argon + 2%Oz

25 l /min Argon 25 l /m in Inco Flux 6

Inco Flux 6 Inco Flux 6 Argon 10 l /min

Argon 10 l /m in

-

80

130 85

130 80

220

230

230 280

210 300 145

140 40 X 10- '

Steel 5, 18%Cr/5%Ni/2.7%Mo

------Argon + 2%0 2

27 l /min Argon 25 l /min Inco flux 6 Inco flux 6 Argon 10 l /m in Argon 10 l /min

-

160 85

145 80

130 75

230

230 210 310 140 140

36 X 10"*

Voltage, V

25

25 25 24

24

26 11

11

11 28

140 x 10:'

24

24 25 25 25 27

22

22 27

31 30 11

11 140 X 103

24 24 25 24 25 25 22

22 31 26 11 11

140 X 10;*

Speed, m m / m i n " "

200

180 165 360

360

430 100

125

95 400

1125

200

200 180 200 200 310

280

280 360

360 360 95

95 1250

200 180 250 180 200 130 280

280 360 400 95

100 1000

Input energy,

k j / m m " "

0.7

0.7 0.7 1.0

1.0

0.8 0.9

1.1

1.0 1.1 0.3

0.6

1.2 0.7 1.2 0.6 1.2

1.2

1.2 1.2

1.1 1.5 1.0

1.0 0.3

1.1 0.7 0.8 0.6 1.0 0.9 1.5

1.1 1.1 1.2 1.0 0.9 0.3

No. of passes"''

10

10 10 7

8

6 20

1.3

14 5 1

5

3 5 3 5 4

3

3 4

3 3

13

I I 1

5 8 6 7 6 8 5

5 3 3

12 14

1

in.) thick plate, grains of up to 0.5 mm (0.02 in.) width and 50 mm (1.97 in.) length were observed.

On welding, considerable HAZ grain growth occurred. Wi th the sheet mate­rial, the peak grain size in the HAZ was about ASTM 2, although that of the weld metal was considerably larger. In the heavier gauge plate, HAZ grain sizes of about ASTM 0 were devel­oped. Some subgrain structure was noted adjacent to the fusion bound­ary, particularly wi th nonmatching composit ion consumables.

Matching composition weld metals showed pronounced epitaxial growth, regardless of process. Ferrite contents of welds made with austenitic stainless steel consumables were quite high. With nickel-base weld metal and all processes, some small hot cracks were found, especially initiating from the fusion boundary. Weld metal cracking was also observed in the matching composition filler SA weld, while HAZ cracking was noted in the GMA weld made using a high nickel consumable, the cracks propagating along large

grains into base metal outside the grain-coarsened region.

Steel 4: Ferrite-Austenite (26%Cr/ 5%Ni/1.5%Mo). This duplex alloy con­tained 20-30% austenite-Fig. 5. The matrix ferrite grain size was quite small (ca ASTM 6), wi th the primary austen­ite banded by rolling. Secondary aus­tenite produced during final solution treatment had formed mainly at ferrite grain boundaries, although not always parallel to the rolling plane.

On heating during welding, trans­formation to ferrite occurred in the

342-s I NOVEMBER 1977

Fig. 7—Strauss test sample geometry, specimen bent as indicated

S t e ,;••'• '*••*. ->

, , . : . . . - . A . v £ £ ? . * - * * * * - • • S * - W . - - . .»».«.-, _ . , - . h— —»— -*_*--__ «5 »^&i ' : ' - - -,-.>.=u— „.

^ "V: :': v. •-•^.sSk

Fig. 2—Microstructures of steel 1: A (top)-base metal, X320; B (middle)-weld metal and high temperature HAZ, X200; C (bottom)-Low temperature HAZ, X200 (50% reductions on reproduction)

HAZ, wi th associated grain growth. On cooling, austenite reformed at grain boundaries and to a lesser extent within the grains. Carbide and nitride

Fig. 3—Microstructure of steel 2: A (top)-base metal, X630; B (bottom)-weld metal and high temperature HAZ, X600 (50% reductions on reproduction)

decoration of subgrain boundaries was evident, wi th a precipitate-depleted area around grain boundaries. As assessed visually, repeated thermal cycling in a multipass weld slightly increased the total austenite content, wi thout affecting the peak grain size.

The matching or near-matching composition weld metals showed a coarse, columnar structure wi th epi­taxial growth from the HAZ. Some intergranular austenite was observed, the austenite being least in the autoge­nous GTA and EB weld metals. Welds made wi th Type 316L filler metals showed increased ferrite contents, wi th austenite and ferrite tending to form in a parallel lath-like structure. Using nickel-base filler metals, a

Fig. 4—Microstructure of steel 3: A (top)-base metal, 13 mm plate, X100; B (middle)—weld metal and HAZ of CTA weld in 13 mm plate with matching compo­sition filler metal, X50; C (bottom)-HAZ cracking in CMA weld in steel 3, using nickel-base filler metal, X80 (50% reduc­tions on reproduction)

sharply defined fusion boundary was apparent, wi th little correspondence between HAZ grain size and weld metal solidification structure.

Wi th all processes, nickel-base weld metals showed small hot cracks adja­cent to the fusion boundary. In addi­t ion, some weld metal solidification cracking was observed in the matching composition SMA and SA welds.

Steel 5: Ferrite-Austenite (18%Cr/ 5%Ni/2.7%Mo). Microstructural obser­vations were essentially the same as for steel 4 with respect to transforma­tion behavior during welding and weld metal characteristics—Fig. 6. The base metal of steel 5 contained more aus­tenite (ca 40%); compared wi th steel 4, the HAZ structures showed finer grain sizes and higher austenite contents, while the transformed HAZ's were narrower. Weld metal cracking was noted wi th nickel-base filler metals.

W E L D I N G RESEARCH SUPPLEME NT I 343-s

Strauss Tests—Tables 6 and 7

Steel 1: AISI 430 (17%Cr). The GTA w e l d e x a m i n e d suf fered severe in ter ­granu lar H A Z attack—Fig. 7. I n t e rg ran ­ular a t tack also occu r red in t h e w e l d me ta l , a l t h o u g h to a lesser ex ten t . A d d i t i o n a l l y , fa i lu re o n b e n d test o c c u r r e d t h r o u g h the w e l d m e t a l , p resumab ly in c o n s e q u e n c e o f t h e par t i cu la r l y coarse w e l d meta l s t ruc­t u r e d e v e l o p e d .

Steel 2: Ferrite-Martensite (IbXCrl 2.7%Ni/l%Mo/0.5%Nb). N o c rack ing or in te rc rys ta l l i ne at tack was o b ­served.

Steel 3: Ferritic ELI (26%Cr/l%Mo). Cons ide r i ng first the n o r m a l Strauss test p r o c e d u r e , t he G T A w e l d in 2.5 m m (0.10 in.) th i ck sheet d i sp layed in te rgranu la r a t tack in t he w e l d meta l and H A Z - F i g . 8. In t he 13 m m (0.5 in.) t h i ck p la te , us ing m a t c h i n g f i l ler meta ls and a rgon sh ie ld i ng , b o t h G T A and G M A w e l d s w e r e resistant to in te rc rys ta l l i ne at tack as was the EB w e l d . The submerged arc w e l d s also appeared resistant to a t tack , a l t h o u g h b o t h w e l d s s h o w e d s o m e br i t t leness. In all o t he r cases, i n c l u d i n g the A r / 1%0_, G M A w e l d w i t h m a t c h i n g f i l ler me ta l , i n te rc rys ta l l i ne p e n e t r a t i o n was obse rved .

Penet ra t ion rates w e r e genera l l y h igher in t he m o d i f i e d Strauss test t h a n w i t h t h e n o r m a l p r o c e d u r e . Results, h o w e v e r , essent ia l ly c o n ­f i rmed those f r o m t h e n o r m a l p r o c e ­du re , a l t h o u g h t h e G M A w e l d w i t h Ar sh ie l d i ng gas and m a t c h i n g f i l ler meta l s h o w e d in te rc rys ta l l i ne w e l d meta l and H A Z at tack , w h i l e m a r k e d c o r r o ­s ion o f t h e S M A N i -base w e l d meta l was obse rved .

Steel 4: Ferrite-Austenite (26%Cr/ 5%Ni/1.5%Mo). The f o l l o w i n g p o i n t s can be m a d e f r o m the n o r m a l Strauss test results:

1. W e l d s m a d e us ing 316L c o n s u m ­ables s h o w l o w suscep t ib i l i t y to in te r ­c rys ta l l ine at tack in H A Z a n d w e l d meta l apart f r o m a u t o g e n o u s GTA roo t runs.

2. A l t h o u g h w e l d meta l d e p o s i t e d us ing n e a r - m a t c h i n g c o m p o s i t i o n consumab les is resistant to in te rc rys­ta l l i ne a t tack , t he associated H A Z is suscept ib le to such c o r r o s i o n .

3. A s imi la r bu t less m a r k e d e f fec t t han obse rva t i on (2) above exists w i t h n icke l -base f i l ler meta ls .

4. EB w e l d s are h igh l y sensi t ive to in te rc rys ta l l i ne at tack in t h e H A Z a n d w e l d meta l .

As w i t h steel 3, h igh rates of w e l d meta l a t tack we re observed in t he m o d i f i e d Strauss test w i t h t he SMA N i -base w e l d s .

Steel 5: Ferrite-Austenite (18%Cr/ 5%Ni/2.7%Mo). A l l w e l d s w e r e resis­tan t t o in te rc rys ta l l i ne at tack as

V

AQM^A

, y - .-' r

—- - - -"V -M-.-r^r;v_S./"='-

•AAity-Fig. 5—Microstructure of steel 4: A (top)-base metal, X400; B (bottom)-weld metal and HAZ of GMA weld using near-matching filler metal, X100 (50% reductions on reproduction)

assessed by t h e no rma l Strauss p r o c e ­du re , apart f r o m the EB w e l d w h i c h suf fered H A Z and w e l d me ta l sens i t i -za t i on . Some t e n d e n c y fo r increased pene t ra t i on in a u t o g e n o u s GTA roo t runs was obse rved .

Severe w e l d meta l co r ros i on was n o t e d o n m o d i f i e d Strauss tes t ing w i t h t he SMA w e l d m a d e us ing N i -base f i l ler meta ls .

Potentiostatic Tests-Tables 8 to 10

A n o d i c p o l a r i z a t i o n curves fo r t h e base meta ls and typ ica l n o n m a t c h i n g c o m p o s i t i o n w e l d meta ls are s h o w n in Fig. 9. These curves i l lus t ra te t he general co r ros ion behav io r o f t h e d i f fe ren t mater ia ls,7 and are g iven p r imar i l y for re fe rence. Regions of a t tack a n d passiv i ty are s h o w n .

It w i l l be n o t e d that unde r m o d e r ­ately o x i d i z i n g c o n d i t i o n s , the c r i t i ca l

Fig. 6-Microstructures of steel 5: A (top)-base metal, X125; B (middle)-weld metal and HAZ of SMA weld using 316L electrodes, X100; C (bottom)—fusion boun­dary crack in SA weld in steel 5 using nickel-base filler metal, X250 (50% reductions on reproduction)

po ten t i a l for a t tack o n any par t i cu la r w e l d area shou ld be as l o w as poss ib le . If, fo r e x a m p l e , the c r i t i ca l p o t e n t i a l for an H A Z is above that fo r t he base me ta l , w e l d i n g must be cons ide red to have had an adverse e f fec t o n t he co r ros ion resistance o f t he m a t e r i a l . ' 7

U n d e r h igh l y o x i d i z i n g c o n d i t i o n s , it

Table 6—Results of Strauss Tests on GTA Welds in Thin Sheet Materials

Steel

No. 1

No. 2

No. 3

Control

Strauss

C S C S c s

Bend test""

F F P P P F

Observation of crack

pos i t ion" "

Weld Weld No cracks No cracks No cracks HAZ and FB

Depth of penetration in mm/year

Weld HAZ

3.8 6.4

0.03 0.06

2.3 2.9

Base metal

0.06

0.03

0.19

Assessment

F

P

F

I-,F—fall; P-pass. ""Weld—weld metal; HAZ—heal affected zone; FB — fusion boundary.

344-s I N O V E M B E R 1 9 7 7

Table 7—Results of Strauss Tests for Plate Samples

Strauss test results Modi f ied Strauss

test results '"

Weld no. Process

SMA

2 SMA

3 SMA

4 GMA argon

GMA 5 argon

+ 1 % 0 2

6 GMA argon

7 GTA

8 GTA

9 GTA

10

11

12

SA

SA

SMA

SMA

SMA

SMA

SMA

SMA

GMA

GMA

Filler metal

Type 309L-MO

Type 316L

Ni-base

Matching composit ion

Matching composit ion

Ni-base

Matching composit ion

Matching composit ion

Ni-base

Matching composit ion

Ni-base

HAZ" " grain Control

Send Observations

of crack (ASTM) Strauss test"" posit ion

Steel 3, 26%Cr/1%Mo.

C P No cracks

Depth of penetration

mm/year

Depth of penetration,

mm/year

Matching composit ion

Matching composit ion

Type 316L

Type 316L

Ni-base

Ni-base

Near matching composit ion

Type 316L

2

2

2

1

> 1

> 1

1 and 2

1

1 and 2

1

1

> 1

1

3

1

1

1

3; 1 in TIG root

1

> 1

S

C

S

C

S

C

S

c s c s c s c s c s c s c s c s

c s c s c s c s c s c s c s c s

F

p

F

F

F

P

P

P

F

F

F

P

P

P

P

F

F

F

F

P

P

P

P

Fusion boundary

No cracks

Fusion boundary

No cracks

Fusion boundary

No cracks

No cracks

No cracks

HAZ and FB

Fusion boundary

Fusion boundary

No cracks

No cracks

No cracks

No cracks

HAZ and FB

HAZ and FB

Weld

Weld

Small tear in FB

Small tear in FB

No cracks

No cracks

Steel 4, 26%Cr/5%Ni/1.5%M

F Small crack defect

F

F

P

F

F

F

P

P

P

P

P

F

F

P

P

Weld

Weld

No cracks

Fusion boundary

Fusion boundary

Weld crack

No cracks

No cracks

No cracks

No cracks

No cracks

HAZ and FB

HAZ and FB

No cracks

No cracks

Weld metal HAZ'

0.03 1.3

Base Assess-metal ment

0.09

Weld Base Assess-metal HAZ" " metal ment " "

0.09

0.03

1.12 0.12

0.95 0.16

0.03 0.09 0.16

0.12 1.2

0.79 0.4

0.16

0.16

0.49 0.32 0.12

0.15 0.32 0.12

0.06 1.7 0.07

0.12 0.38 0.09

0.2 0.32 0.16

0.09 0.06 0.06

0.06 0.12 0.06

0.03 (1.12 0.32 0.06

(b)) 0.09 (1.74 0.29

(b)) 0.09 (0.47 0.22

(b»

0.1

0.09

0.06 0.15 0.06

0.03 0.03 0.03

0.06

0.03

1.3

0.12

0.06

0.03

0.09 • 1.8 0.63 F

(1.4)""

0.18 2.3 0.36 F

18.8 0.1 0.27 F

2.28 1.62 0.38 F

ND ND ND ND

1.35 0.95 0.9 F

0.1 0.95 0.95 P

0.09 0.47 0.57 P

ND ND ND ND

ND ND ND ND

0.19 0.45 0.27 P

0.18 0.09 1.2 P

0.63 (1.08)""

2.0 (5.2)""

0.18 (1.4)""

0.18 (0.9)""

2.07

0.72

0.57

1.99

1.35

1.17

0.27

0.09

0.27

1.62

0.54

P

P

F

F

F

15.93 0.54 0.1c

0.26 2.19 1.42

0.28 0.72 0.26 P

(Cont inued on next page)

is des i rab le tha t t he c r i t i ca l p o t e n t i a l for b r e a k d o w n o f passiv i ty be as h igh as poss ib le . Few service m e d i a o the r

t han n i t r i c ac id are su f f i c i en t l y h igh l y o x i d i z i n g fo r co r ros ion to o c c u r at t he " t o p " of t he passive range. Preferent ia l

in te rc rys ta l l i ne at tack in p rac t i ce is p r e d o m i n a n t l y associated w i t h the mode ra te l y o x i d i z i n g reg ime, and at-

WELDING RESEARCH SUPPLEMENT I 345-s

Table 7—Results of Strauss 1

Weld Filler no. Process metal

rests for Plate

HAZ"" grain size

(ASTM)

Samples—Contini

Control Rend

Strauss test11"

led

Strauss test

Observations of crack posit ion

results

Depth of penetration

mm/year

Weld Base metal HAZ'* ' metal

Assess­ment

Modi f ied Strauss test results"'1

Depth of penetration,

mm/year

Weld Base Assess-metal HAZ"*' metal ment " "

P No cracks

10

11

12

13

14

IS

GMA

SA

SA

SA

GTA

GTA

I B

SMA

SMA

SMA

SMA

SMA

SMA

ro

11

12

13

SA

SA

GTA

GTA

Ni-base

Near matching composit ion

Type 316L

Ni-base

Near matching composit ion

Ni-base

3 & 1

3 & 1

Type 309L-MO

Type 309L-MO

Type 3I6L

Type 316L

Ni-base

Ni-base

GMA Type 316L

GMA Ni-base

Type 316L

Ni-base

Type 316L

Ni-base

3 & 4

2 & 3

s c S

c S

c s c s c s c s

c s c s ( s c s c s c s c s c s c s c s c s c s c s

F

I

F

F

F

F

F

P

F

F

F

F

F

Stee

P

P

P

P

F

P

F

F

P

P

P

P

P

P

P

P

P

P

P

P

P

P

P

P

P

F

HAZ

HAZ

HAZ

Fusion boundary

Fusion boundary

Weld

Fusion boundary

No cracks

HAZ and FB

Fusion boundary

Fusion boundary

Weld

Weld

:l 5, 18%Cr/5%Ni/2.7%M

No cracks

No cracks

No cracks

No cracks

Weld

No cracks

We ld

Weld

No cracks

No cracks

No cracks

No cracks

No cracks

No cracks

No cracks

No cracks

No cracks

No cracks

Weld

No cracks

No cracks

No cracks

No cracks

No cracks

No cracks

HAZ

0.15 0.75 0.06 F 0.29 0.46 0.76 F

0.06 1.9 0.06 F ND ND ND ND

0.03 0.22 0.06 P 0.36 0.09 0.18 P

0.03 0.94 0.03 F 1.53 1.35 0.45 F

1.2 0.9 0.06 F 0.38 0.19 0.29 F

0.09 0.12 0.06 P ND ND ND ND

14.2 9.6 0.06 F ND ND ND ND

0.6 (0.47 (b)) 0.3 (0.48 (b)) 0.16 (0.24 (b))

0.24

0.16 (0.3 (b)) 0.06 (0.23 (b))

0.03

0.09

0.03

0.15

0.16

0.19

2.2

0.06

0.3

0.09

0.24

0.3

0.13

0.32

0.16

0.25

0.16

0.22

0.12

0.63

0.06

0.09

0.06

0.03

0.12

0.06

0.09

0.19

0.09

0.09

0.06

0.03

0.06

P

P

P

P

P

P

P

P

P

P

P

P

F

0.09 (3.2)"" 0.27 0.27 P

0.19 (1.4)"" 0.19 0.67 P

0.09 (0.9)"" 0.36 0.72 P

0.09 (0.19)"" 0.19 0.54 P

6.7 0.19 0.38 F

24.5 0.67 0.48 F

0.48 0.48 0.38 P

0.18 0.27 0.27 P

0.9 0.18 0.63 F

0.09 0.9 0.63 F

0.45 0.18 0.63 P

ND ND ND ND

1.35 0.27 0.54 F

•'HAZ-heat-affected zone. tt'P-pass; F-fail. "ND-not determined. "'Penetration for autogenous GTA root.

346-s I NOVEMBER 1977

Table 8—Potentials for Intergranular Attack in Mildly Oxidizing Conditions for GTA Runs in Thin Sheet Samples—Standard Electrolyte

"5a--.i:- *® Fig. 7—Intercrystalline attack in CTA weld in steel alter Strauss testing, X75 (50% reduction on reproduction)

Fig. 8-A (top)-cracking in GTA weld in steel 3 after Strauss testing and bending, X5; B (bottom)-intergranular penetration in HAZ of 309L-Mo SMA weld in steel 3, X320 (50% reductions on reproduction)

tention is concentrated on such condi­tions in the present report.

Two points should be made: 1. In the AISI 430 base metal,

marked attack occurred around car­bide particles at potentials above that corresponding to breakdown of the passive f i lm at grain boundaries, or over the entire material. This attack is associated wi th the kink in the polari­zation curve at above 4-0.3 V—Fig. 9. Such attack was not observed in other alloys and, apart from where indicated in Tables 8-10, all potentials corres­pond to the observed init iation of intercrystalline attack.

2. Increasing the exposure t ime from 20 min to 16 or 64 h resulted in a general rise in the critical potentials recorded, presumably due to a kinetic effect of sluggish passive fi lm break­down in the alloys examined.

The addition of 3% NaCI to the standard electrolyte has the effect of widening the active loop, increasing the peak current density,9 and raising

Steel

No. 1

No. 2

No. 3

Table 9-Potcn

Steel

No. 3

No. 4

No. 5

Time of exposure

20 min

16 h

20 min 16 h

20 min 16 h 64 h

tial for Intergranu

Time of exposure

20 min 16 h 64 h 20 min 16 h 20 min 16 h

Base metal,

v„ + 0.27

(around carbides) + 0.1

(grain boundaries) + 0.38

(around carbides) + 0.14

(grain boundaries) + 0.14 + 0.27 -0.02 -0.02 + 0.09

lar Atfack in Base Metal

Moderately oxidizing, V,,

+ 0.03 + 0.03 + 0.09 + 0.09 + 0.18 + 0.08 + 0.09

HAZ, V„

+ 0.18

+ 0.39 + 0.09 + 0.20 -0.02

0 + 0.11

Plate Samples

Highly oxidizing, V n

+ 1.33

-—

+ 1.32

-+ 1.30

—

Weld metal,

v„

+ 0.18

+ 0.39 + 0.09 + 0.20 -0.04

0 + 0.1

3% NaCI added, VH

+ 0.02 + 0.05

-+ 0.24 + 0.29 + 0.08 + 0.09

the potential below which intercrys­talline attack is observed. However, the chloride addit ion has had no effect on either the observed incidence of preferential HAZ or weld metal corro­sion relative to weld metal, or on the apparent rate of passive fi lm break­down.

Steel 7; A/5/ 430 (17%Cr). From the results in Table 8, the welding has raised the potential below which intercrystalline attack occurs in steel No. 1 by 0.25 V. This is a considerable increase, and indicates that the mate­rial wi l l be potentially susceptible to preferential corrosion in the HAZ and weld metal in service. Further, the critical potential for intercrystalline attack is higher than for any other alloy. Thus, the risk of service attack is greatest for AISI 430 material in that a wider range of environments wi l l be expected to cause HAZ corrosion.

Steel 2: Ferrite-Martensite (16%Cr/ 2.5%Ni/17oMo/0.5%Nb). It was appar­ent that welding had little effect. In fact, the results indicate welding to improve corrosion resistance in that the potential for attack on the weld metal and HAZ was below that for base metal.

Steel 3: Ferritic ELI (2670Cr/17oMo). All HAZ's showed similar critical potentials to base metal, and it wi l l be noted that these potentials are lower than for the other steels. Nickel alloy weld metals showed attack at poten­tials where the base metal steel

remained passive—Fig. 10. This was also the case wi th the Type 309L-Mo and 316L filler metals; in the latter welds, the autogenous GTA root runs also showed critical potentials higher than for base metal. Enhanced fusion boundary corrosion was noted wi th nickel-base filler metals—Fig. 11.

5tee/ 4: Ferrite-Austenite (26%G7 5%Ni/1.5%Mo). Close to the active loop, HAZ intercrystalline attack was observed at potentials slightly above the passive fi lm breakdown potential recorded for the base metal. No partic­ular trend in critical potential for the HAZ could be discerned wi th welding conditions, apart from the potential being lowest wi th SMA welding and highest for the EB weld.

Under highly oxidizing conditions, HAZ and base metal corrosion behav­ior seemed similar, apart from greater susceptibility to attack in the EB weld. At potentials causing corrosion, the HAZ's of welds made wi th the near-matching filler metal showed marked intercrystalline attack close to the fusion boundary—Fig. 12.

In the base metal and in low temper­ature HAZ's, the ferrite/austenite in­terphase boundaries were more heavi­ly attacked than the ferrite grain boundaries.

With the near-matching and Type 316L consumables, similar passive/ active transition potentials to those of the HAZ's were observed. This was not the case for the nickel alloy weld

W E L D I N G RESEARCH SUPPLEMENT I 347-s

Table 10—Potentials for Intergranular Attack in Plate Samples

Weld no.

1

2

3

4

5

6

7

8

9

10

11 12

Process

SMA

SMA

SMA

GMA argon GMA

A + 1 % 0 . GMA argon GTA

GTA

GTA

SA

SA EB

Filler metal

Type 309L-MO

Type 316L

Ni-base

Matching composit ion

Matching composit ion

Ni-base

Matching composit ion

Matching composit ion

Ni-base

Matching composit ion

Ni-base

SMA

2

3

4

5

6

7

SMA

SMA

SMA

SMA

SMA

GMA

Matching composit ion

Type316L

Type 316L

Ni-base

Ni-base

Matching composit ion

Time of exposure HAZ"

Steel 3, 26%Cr/1%Mo

20 min +0.04 16 h +0.04

20 min +0.03 16 h +0.03

20 min +0.04" 16 h +0.04 r

20 min -0 .1 16 h +0.03

20 min +0.03 16 h ND

20 min -0.02 16 h +0.02'

20 min 0.01 16 h 0

20 min 0 16 h 0.04 64 h 0.06

20 min -0.02 ND

20 min 0.02 16 h 0.06

20 min 0.01c' 20 min -0.03

Steel 4, 26%Cr/5%Ni/1.5%Mo

Matching composit ion 20 min +0.13 16 h +0.19

20 min +0.11 16 h +0.16

20 min +0.14 16 h +0.22

20 min +0.13 16 h +0.19

20 min +0.12 16 h +0.19

20 min +0.11 16 h +0.21

20 min +0.13 16 h +0.21

Standard electrolyte V „ 3% NaCI added, V„

WM« HAZ" W M "

+ 0.13b

+ 0.15 + 0.14"

0.18" 0.38e

0.38'' 0

+ 0.05 + 0.05 ND

+ 0.31 + 0.36

0.01 0 0.01 0.04 0.06

+ 0.28 ND 0.02 0.01

+ 0.30 -0.03

+ 0.12 + 0.19

+ 0.11 + 0.20 + 0.12 + 0.21

+ 0.13 + 0.19 + 0.39 + 0.40 + 0.39 + 0.38 + 0.13 + 0.20

+ 0.05 ND

+ 0.08 ND

+ 0.04 ND

+ 0.03 ND

+ 0.01 + 0.05 + 0.02 + 0.05

0 + 0.06 + 0.03 ND ND

+ 0.08 + 0.05 + 0.05 ND

+ 0.08 + 0.02

+ 0.18 + 0.25

+ 0.19 + 0.26 + 0.32 + 0.39

+ 0.29 + 0.38 + 0.26 ND

+ 0.27 ND

+ 0.23 + 0.33

(Cont inued on

•I-0.20 ND

+ 0.23 ND

+ 0.56 ND

+ 0.03 ND

+ 0.02 + 0.05 + 0.43 + 0.50

0 + 0.06 + 0.03 ND ND

+ 0.42 + 0.47 + 0.06

ND + 0.45 + 0.02

+ 0.19 + 0.25

(0.27)"" + 0.18 + 0.25 + 0.32 + 0.38

( + 0.40)"" + 0.30 + 0.32 + 0.56 + 0.60 + 0.56 + 0.62 + 0.22 + 0.31

next page)

meta ls , w h i c h we re cons i s ten t l y at ­tacked at a b o u t 200 m V h igher t h a n t h e o thers . There was some p r o p e n s i t y f o r m o r e rap id at tack t o o c c u r at t he fus ion bounda r ies o f w e l d s m a d e using n o n m a t c h i n g c o n s u m a b l e s than e l sewhere in t he samples e x a m i n e d and fo r a u t o g e n o u s GTA roo t runs also to s h o w an increase in co r ros i on rate.

5tee/ 5; Ferrite-Austenite (78%Cr/ 57oNi/2.77oMo). Observations made are largely similar to those for steel 4. Under moderately oxidizing condi­tions, the critical potentials for attack on the HAZ's, or stainless steel weld metals, were above the potential obtained for base metals, wi th autoge­nous GTA root runs tending to corrode more rapidly.

The EB weld showed greatest sus­ceptibil ity to attack. Again, the corro­sion characteristics of nickel alloy weld metals differ from those of the base metal. Ferrite/austenite interface attack was less marked than wi th steel 4, wi th attack on the austenite phase being more apparent.

Discussion

Welding Characteristics

Overall, the base metals studied are similar to austenitic stainless steels with regard to carrying out the welding operation per se. Some lack of fusion defects was observed in the test welds, especially at low heat inputs. However, the fusion and penetration characteristics of the materials exam­ined were comparable to those of austenitic stainless alloys. Conse­quently, conventional stainless steel welding techniques can be considered to be directly applicable to the present alloys.

From the welding viewpoint, the greatest difference between the pres­ent materials and common austenitic grades is the increased risk of weld metal cracking in the former case. Cracking was found wi th both match­ing and nickel-base filler metals, the morphology generally indicating the cracks to have formed during solidif i­cation. In nickel-alloy weld deposits,

cracks apparently initiating at the fusion boundary were frequently ob­served.

Typical crack lengths were of the order 0.5-1.0 mm (0.02 to 0.04 in.). While the cracks may not be of prac­tical significance, this observation does mean that careful consideration must be given to the use of nickel alloy filler metals for the stress corrosion resistant stainless steels examined. There was no evidence that such cracking was dependent on welding conditions. Also, it would appear to stem from the formation of a region having a crack-sensitive composit ion as a result of di lut ion of the filler metal by a low-nickel base metal.

In the fully ferritic steel 3, weld metal cracks in submerged arc welds made using matching composit ion filler metals appeared to have formed by cleavage at low temperature. It is well established that the toughness of fully ferritic stainless steels is highly dependent on material purity/1 and some pick-up of interstitial elements

348-sl N O V E M B E R 1977

Table 10-Potentials for Intergranular Attack in Plate Samples-Continued

Weld no.

8

9

10

11

12

13

14

15

1

2

3

4

5

6

7

8 9

10

11

12

13

Process

GMA

GMA

SA

SA

SA

GTA

GTA

EB

SMA

SMA

SMA

SMA

SMA

SMA

GMA

GMA SA

SA

GTA

GTA

EB

Filler metal

Type 316L

Ni-base

Matching composi t ion

Type 316L

Ni-base

Matching composi t ion

Ni-base

Steel 5,

Type 309L-Mo

Type 309L-MO

Type 316L

Type 316L

Ni-base

Ni-base

Type 316L

Ni-base Type 316L

Ni-base

Type 316L

Ni-base

—

Time of exposure

20 min 16 h

20 min 16 h

20 min 16 h

20 min 16 h

20 min 16 h

20 min 16 h

20 min 16 h

20 min

18%Cr/5%Ni/2.7%Mo

20 min 16 h

20 min 16 h

20 min 16 h

20 min

20 min 16 h

20 min 16 h

20 min 16 h

20 min 20 min

16 h 20 min

16 h 20 min

16 h 20 min

16 h 20 min

16 h

M d l l U d t U eifc

HAZ"

+ 0.12 + 0.28 + 0.12 + 0.29 + 0.14 + 0.35 + 0.16 + 0.23 + 0.12 + 0.27 + 0.09 + 0.25 + 0.09 + 0.35 + 0.26

+ 0.13 0.13

+ 0.14 + 0.18 + 0.13 + 0.16 + 0.13 + 0.18 + 0.18 ND

+ 0.14 + 0.15 + 0.12 + 0.19 + 0.14 + 0.12 + 0.20 + 0.12

ND + 0.10 N D

+ 0.12 ND

+ 0.30 + 0.45

cuoiy ie v H W M "

+ 0.14 + 0.19 + 0.31 + 0.30 + 0.14 + 0.19 + 0.18 + 0.23 + 0.35 + 0.30 + 0.09 + 0.30 + 0.32 + 0.40 + 0.65

+ 0.14 0.14

+ 0.17 + 0.19 + 0.16 + 0.17 + 0.17 + 0.17 + 0.38 + 0.38 + 0.37 + 0.39 + 0.17 + 0.20 + 0.36 + 0.17 + 0.17 + 0.29 ND

+ 0.12 N D

+ 0.27 ND

+ 0.60 + 0.63

J/O I N d C I <J

HAZ"

+ 0.31 + 0.34 + 0.23 ND

+ 0.22 + 0.33 + 0.30 + 0.31 + 0.23 + 0.34 + 0.16 + 0.28 + 0.27 + 0.33 + 0.26

+ 0.22 + 0.28 + 0.22 + 0.27 + 0.25 + 0.31 + 0.26 + 0.31 + 0.28 ND

+ 0.26 ND

+ 0.20 + 0.25 + 0.28 + 0.22 + 0.26 + 0.30 + 0.26 + 0.21 + 0.24 + 0.33 + 0.26 + 0.38 + 0.50

uueu, v || W M "

+ 0.30 + 0.33 + 0.54 + 0.55 + 0.24 + 0.31 + 0.33 + 0.33 + 0.59 + 0.69 + 0.18 + 0.30 + 0.44 + 0.45 + 0.45

+ 0.22 + 0.28 + 0.22 + 0.27 + 0.25 + 0.32 + 0.23 + 0.31 + 0.56 ND

+ 0.56 + 0.63 + 0.23 + 0.27 + 0.42 + 0.24 + 0.28 + 0.44 + 0.48 + 0.24 + 0.27 + 0.48 + 0.51 + 0.38 + 0.58

"HAZ-heat-affected zone; WM-weld metal "Attack below this potential on both autogenous GTA and bulk weld metal. •Attack most apparent at fusion boundary.

"Critical potential for GTA root runs.

during submerged arc welding must be anticipated, with consequent de­creased capacity for accepting welding strains wi thout cracking. Such an effect of contamination is consistent wi th results of other workers. Unless special consumables can be devel­oped, the submerged arc process cannot be recommended for this material.

Steel 3 also showed HAZ cracking. As in the submerged arc weld metal the cracks appeared brittle, and to have formed at low temperature.

Evaluation of Corrosion Test Methods

Susceptibility to intercrystalline at­tack was found wi th the Strauss test. In principle, the method appears as applicable to the present alloys as it is to austenitic steels. In a number of instances, test welds showed poor bend ductil ity. However, discrimina­

t ion between different weld areas was successfully obtained by the use of metallographic examination to define the depth of penetration during expo­sure to the acid copper sulphate. At the same time, it may be noted that intergranular attack in the Strauss test electrolyte takes place only in areas heated too close to the solidus temperature during welding. Hence, attack is extremely localized and, in an HAZ, may arise over only a few grain diameters from the fusion boundary. Careful examination is necessary to locate and measure such intergranular penetration. Moreover, the procedure cannot be regarded as a straightfor­ward standard test method, as is the normal procedure called for in BS 1449: Part 4. 1967.

In ascribing a "pass" or " fa i l " assess­ment to the welds tested by the Strauss method, consideration is given both to the observed bend behavior, and to

the rates of penetration recorded. It wi l l be noted that, especially with the modif ied procedure, appreciable at­tack could take place on base metal and that the rate of such attack showed considerable variation—Table 7. This has to be borne in mind when deciding whether or not a welding operation has significantly increased the risk of intercrystalline corrosion.

The practical application of data generated from a standard test proce­dure such as the Strauss test wi l l depend on the view taken of the rele­vance of the test to anticipated service conditions. In this regard, reference wi l l probably be made to whether or not a history exists of intercrystalline attack in welded austenitic alloys.

In principle, it is advantageous to employ a test method capable of giving information pertinent to as wide a range of environments as possi­ble. However, it is clear from Tables

W E L D I N G RESEARCH SUPPLEMENT I 349-s

!

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I \{/\ J \ \,

"^x* -—— ^'yA-'^-''A

"._ •*

oe

o-a

0 <

0 3

0 2

0 1

0

- 0 ,

. i i i

11 'ML

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i 1 Jl

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• •

• Steel 1 • Steel 2

17% Cr 16'/.Cr 2-7'ANi

Steel i 26 "/. Cr 5/V; K V.Mo Steel 5 18 '/.CrS'/.Ni

2-7V.M0

-^"—-~i>—~

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i in..,.! ! IH....I : ' ...I i III.„J u n ,-J 2 J 4! ,„-2 2 1 «5 ,„•> J 3 45 JO ' 3 t $ to2

Current density , mA/cm2

Fig. 9—A (top left)—typical potentiostat scans of base metal steels; B (top right)—typical potentiostat scans of representative weld metals; C (bottom right)—typical potentiostat scans of base metal steels in the presence of 3% NaCI

1

1

/

Tronspa ss i ve region

1 111

^ < ^ " ___ — • " _ ' _ C ^ - ^ " '

-̂ V" i ( In i

1 1

Highly oxidising

l l l l

""

i- 0-5

0-3

0-1

-0-2

St a

S,

i i ! 1

11 MM

i \

'

l l l l ! 1 '

* -——**

—^s-^-:=i.^r._ Active loop C— ,.

r*J ' A 2

1 1 1 HI Steel. • 4 Ni ba

1

se alloy

i n i " veld metal

• 5 Type 316L MIG weld metal

J Moderate oxidism

ty 7

Current density, mA/cm2

1 111

/

1 i, nil

„ - - - '

= = ~ ^

11 n

' - ' ^ ^ i |

m i

OS

OS

ot

0-3

02

0-1

0

0-1

11

/I

:i;:^>.-

l l l l 1 : 1 1 1 II Steel

I I 11

5 IS V.Cr 5 * 2-7Mo 3 26'/.Cr t'/.Mo

^̂ "̂ *

— ̂

1

i ^ ^ . ̂

III....I ' •,-* ' .

1

Current density. mA/cm2

8-10 that adequate time for passive f i lm breakdown must be allowed in potentiostatic testing to obtain an accurate assessment of material be­havior. This has been pointed out previously."1 In other w o r k " at The Welding Institute on austenitic steels, this has not caused any practical diff i­culties. Nonetheless, it is obvious that

the kinetics of passive fi lm breakdown on the present materials must be closely defined before a definitive potentiostatic test procedure can be established.

From the present results, it is clear that microstructural changes occurring during welding affect corrosion be­havior, and susceptibility to preferen­

tial HAZ attack was observed using the potentiostat. However, two points must be made:

1. The potentiostat indicated there to be little difference between the HAZ's arising wi th the various arc processes studied, yet w i th steels 3 and 4, the Strauss test definitely showed

350-s I NOVEMBER 1977

Fig. 10—Preferential attack on nickel-base SMA weld metal following exposure to + 0.2V,, for 76 hours: steel 3, 13 mm (0.5 in.) thick base metal and autogenous GTA root run, unattacked, X5 (50% reduction on reproduction)

welding procedure to influence sus­ceptibil ity to intercrystalline attack.

2. Wi th nickel-base weld metals, the critical potentials for intercrystal­line HAZ attack lay wi th in the poten­tial range, causing weld metal corro­sion; hence, under moderately oxi­dizing conditions, environments caus­ing HAZ penetration should also cause weld metal attack.

This was not the case, since negli­gible weld metal attack was found on Strauss test samples which were susceptible to HAZ corrosion (al­though weld metal corrosion was observed wi th the modif ied Strauss procedure). It must, therefore, be concluded that the potentiostatic technique for evaluating sensitivity to intergranular attack as developed for plain 18%Cr/10%Ni type austenitic steels7 is insufficiently sensitive to give a definit ive assessment of suscepti­bility with the present steels.

At least wi th steels 3 and 4 the critical potentials for base metal or HAZ attack close to the active loop were higher wi th longer exposure. This was not the case with nickel-base weld metals, which showed consistent breakdown potentials. This latter be­havior is similar to that found previously for austenitic steels.7 Thus, there seems no reason to doubt the applicability of the potentiostatic test procedure to nickel-base weld metals.

The relative merits of the poten­tiostat and Strauss test have been considered elsewhere.'" Essentially, they measure slightly different param­eters. The potentiostat determines whether or not susceptibility to inter­crystalline corrosion exists, and if so, over what potential range such corro­sion can occur. On the other hand, the Strauss tests gives an indication of the rate of attack associated wi th any susceptibility that may exist. For the present, particular attention should be paid to the Strauss test, pending devel­opment of a potentiostatic test proce­dure in which full allowance is made for passive f i lm breakdown.

Fig. 11—Preferential fusion boundary attack in weld in Fig. 10, following exposure for 20 min at + 0.04V,,, X200 (50% reduction on reproduction)

Comparison of Materials

The salient feature of the data in Tables 8-10 is that steels 2-5 suffer less susceptibility to intergranular attack on welding than does steel 1. This was found by both the potentiostat and Strauss tests, and indicates that these recently developed alloys wil l be less sensitive to local loss of corrosion resistance when used in welded fabri­cations than wi l l the traditional AISI 430 material. Steels 2-5 wil l also show advantages in improved general corro­sion resistance over AISI 430, although differences between them must be anticipated in some service media.

From the corrosion tests carried out, three observations in particular can be made:

1. Under the conditions of tests employed, corrosion resistance is least affected by welding wi th steels 2 and 5. No appreciable sensitization to intercrystalline attack was observed in the arc welds studied, only the elec­tron beam weld in steel 5 failing the Strauss test.

2. Both steels 3 and 4 can be welded wi thout inducing susceptibility to preferential intergranular attack in the HAZ or weld metal, but attention to welding procedure is required.

3. The corrosion resistance of the nickel-base weld metals studied is not identical to that of the base metals. These weld metals have narrower passive ranges; they can be expected to suffer preferential attack in some environmental conditions, especially those corresponding to potentials close to the active loop.

Steel 2 is stablized with niobium, and it would seem that this approach is effective in avoiding intercrystalline attack, at least in the thin section size examined. It is probable that intercrys­talline corrosion in the materials exam­ined arises as a result of chromium carbide (and nitride) precipitation and local chromium depletion, in an analo­gous manner to that in austenitic stainless steels."1,1 Indeed, a number of workers have studied the effect of traditional stabilizing elements.'°' ia

The present data support this view

Fig. 12-HAZ of CTA weld in steel 4 with near-matching filler metal showing precipi­tation and attack at grain boundaries, X320 (50% reduction on reproduction)

insofar as the potential range over which intercrystalline attack was ob­served—and the range corresponds to that in which attack on low chromium regions wil l occur. On this basis, the relatively good behavior of steels 4 and 5 can be directly ascribed to the rela­tively large ferrite/austenite inter­phase boundary area; chromium car­bide and nitride precipitation wil l occur at the interphase boundaries as well as at grain boundaries and the density of precipitation and resultant chromium depletion wil l be reduced.

Susceptibility to intergranular attack in steels 4 and 5 was associated wi th a coarse predominantly ferritic struc­ture. The greater tendency of steel 4 to transform to ferrite during the welding cycle is presumably largely responsible for the greater sensitivity of this alloy to a welding operation.

In this context, intercrystalline at­tack has been reported in steel 5 (and to a lesser extent in steel 4) under saline conditions." Subsequent to such attack being observed, more nitrogen has been added to steel 5 to reduce its ferrite-forming tendency. From the present results on current production analysis material, this compositional modification has been successful in avoiding susceptibility to intercrystal­line attack, at least in multipass arc welds in the section size examined.

In steel 4, satisfactory joint corrosion resistance was obtained wi th the SMA and GMA processes and a Type 316L stainless steel filler metal. However, the use of a near matching filler metal consistently gave reduced HAZ corro­sion resistance. It has been noted else­where that the choice of filler metal can influence HAZ corrosion resistance in stress corrosion resistant stainless alloys.1" The effect apparently arises because of diffusion of deleter­ious elements, especially interstitial elements, from the weid metal to a predominantly ferritic HAZ during the welding operation. On this basis, the problem can be ameliorated in steel 4 by attention to consumable composi­t ion. It may be possible to improve

WELDING RESEARCH SUPPLEMENT I 351-s

matters by promoting austenite forma­tion in the HAZ, either by control l ing the cooling rate or by using a postweld heat treatment.

The tendency to form ferrite in steel 4 was highest in autogenous GTA root runs and on EB welding; increased susceptibility to intercrystalline attack was observed in both these situations. Electron beam welding would seem to be of l imited application to this alloy, unless postweld heat treatment can be applied. Wi th autogenous GTA runs it would be recommended that a sealing run be employed, or that the root run should not be exposed to aggressive conditions.

The cause of failure of the GTA weld in thin gauge steel 3 is not clear, although an earlier footnote regarding material composition should be noted. There is ample evidence from the results of other recent investigations on steel 3 that stringent control of welding procedure is essential to avoid any contamination.3141" Hence, most satisfactory results were presently ob­tained wi th 13 mm (0.5 in.) thick plate using gas shielded arc or electron beam welding, with matching compo­sition filler metal in the former case.

Other workers have found failure on Strauss testing welds in this alloy made using Type 316 stainless steel and nickel-base filler metals in material thicknesses of several millimeters. There seems litt le doubt that the alloy is highly sensitive to HAZ "contamina­t ion" from nonmatching composit ion weld metal, and control of contamina­tion rather than the final microstruc­ture would appear to be the more important factor in maintaining corro­sion resistance on welding.

Practical Implications

The steels examined in this paper were developed to have corrosion resistance equivalent to that of com­mon austenitic grades. Especially wi th steel 3, it is likely that each wi l l offer specific advantages'in resisting attack in some media.17 However, in the first instance, they wil l probably be speci­fied for applications where the stress corrosion resistance of austenitic al­loys has proven inadequate. Particu­larly wi th the duplex materials, austen­itic stainless steel filler metals have commonly been recommended on the basis of weldment corrosion and mechanical properties.2-'s

Dilut ion from base metals increases the ferrite contpnt typically to about 25%.18 However, because such weld metal contains austenite as the major phase, it is possible that a continuous path for SCC propagation may exist through the deposit. Hence, it has been considered desirable to employ nickel alloy filler metals, on the basis

that the high nickel content wi l l render the weld metal immune to chloride-induced SCC.

This use of nickel-base weld metal must be questioned on the basis of two sets of reasons:

1. The weld metal corrosion resistance is not necessarily equivalent to that of base metal; in many situa­tions, this wi l l not matter since both weld and base metal wi l l be sensibly passive. However, a risk of preferential weld metal corrosion in some media must be accepted. The use of molyb­denum-containing nickel consuma­bles may improve matters, in view of the effect of this element in promoting passivity under moderately oxidizing conditions.1"

2. The use of nickel-base weld metals increases the risk of weld metal cracking, and in this regard the nickel-base weld metals do not confer immunity from SCC.

It has been found that, because of di lut ion, a zone can be formed at the fusion boundary which is fully austen­itic. However, this zone contains substantially less than the 40% nickel generally considered necessary to avoid SCC, and tests have shown that SCC along this zone can occur.16-20

At the present state of the overall test program, it is premature to make definitive welding recommendations for the alloys under investigation. However, on the basis of the corrosion tests carried out, the fol lowing com­ments can be made: Each steel can be welded to give satisfactory results, but it would appear that steel 5 is least affected by a welding operation. The near-matching filler metal used for steel 4 was not employed wi th steel 5 so that the effect of this particular consumable on the HAZ behavior of the latter alloy was not established. Nonetheless, nonmatching consuma­bles were employed wi th no apparent effect, indicating that steel 5 offers considerable flexibility regarding choice of welding procedure.

The single pass EB weld showed some loss of properties in the section size studied. A similar effect may occur in principle wi th arc welding if a mul t i ­pass procedure is not employed to give structural refinement in the HAZ. Thus, it wi l l be necessary in future work to evaluate the effect of higher input energy welding conditions as may be used in practice to obtain increased deposition rates.

The results obtained indicate that steel 4 wi l l be most satisfactorily welded using austenitic stainless steel filler metals and any multipass arc process. However, attention must be given to the possible adverse behavior of an autogenous GTA root run. It has

been demonstrated by other workers that steel 3 can be welded in thick­nesses of up to 12 mm (0.47 in.), wi th satisfactory results.3151"-21 However, it is clear that stringent attention must be paid to avoiding contamination during welding, and, indeed, a close parallel can be drawn between steel 3 and t i tanium. Thus, gas shielded arc or electron beam welding must be regarded as essential for this alloy.

At the present t ime, the commercial future of steel 2 is uncertain. However, accepting that only thin section mate­rial was studied, the high resistance to intercrystalline attack shown after welding is noteworthy. Consequently, there would appear to be scope for development of stress corrosion resis­tant alloys based on a ferrite-marten-site structure.

Conclusion

A study was carried out on recently developed stress corrosion resistant stainless steels to examine the effect of welding on intergranular corrosion resistance. This was done wi th partic­ular reference to moderately oxidizing environmental conditions. The mate­rials examined were AISI 430 (steel 1, included for comparative purposes), a ferritic-martensitic steel (steel 2), an extra low interstitial fully ferritic alloy (steel 3), and two ferritic-austen-itic: materials (steels 4 and 5). The fo l ­lowing conclusions were reached:

1. All the new alloys examined showed advantages over conventional AISI 430 material wi th respect to corro­sion resistance after welding.

2. Steel 2 containing 16%Cr/2.5%Ni/ 1%Mo/0.5%Nb showed no suscepti­bility to intergranular corrosion after GTA welding, although only thin gauge sheet was examined.

3. Opt imum results on steel 3 con­taining 26%Cr/1 %Mo were obtained with gas shielded arc or electron beam welding, where contamination of the weld area was minimal. In both steels 3 and 4, choice of consumable could influence HAZ behavior.

4. The behavior of steel 4 containing 26%Cr/5%Ni/1.5%Mo was broadly comparable wi th that of steel 5. Some reduction in bend ducti l i ty and corro­sion resistance was noted in a number of welds examined, but satisfactory results were obtained using the SMA and GMA processes wi th Type 316L stainless steel filler metals.

5. Steel 5 containing 18%Cr/5%Ni/ 2.7%Mo was little affected by welding under the conditions employed. Cor­rosion resistance similar to that of base metal was obtained using Type 316L or 309L-Mo stainless steel filler metals and a range of arc processes. Electron beam welding caused some loss of

352-sl NOVEMBER 1977

b e n d d u c t i l i t y , and i n d u c e d suscep t i ­b i l i t y to in te rc rys ta l l i ne at tack.

6. The co r ros i on behav io r o f the n icke l a l loy w e l d meta ls s t u d i e d was no t i den t i ca l t o tha t o f t he base meta ls . In many e n v i r o n m e n t a l c o n d i t i o n s t he d i f f e rence w i l l no t be s ign i f i can t , bu t a risk o f p re fe ren t ia l w e l d m e t a l c o r r o ­sion in m o d e r a t e l y o x i d i z i n g m e d i a w i t h n icke l a l l oy f i l le r meta ls mus t be cons ide red .

7. W e l d meta l c rack ing was f o u n d w i t h steels 3-5. This was par t i cu la r l y n o t e d w i t h n i cke l -base f i l ler meta ls , i n i t i a t i ng at the fus ion b o u n d a r y . In a d d i t i o n , steel 3 s h o w e d some suscep­t i b i l i t y to b r i t t l e H A Z c rack ing .

Ac/cnow/edgements

The au thors t h a n k the i r co l leagues at The W e l d i n g Ins t i t u te fo r the i r adv ice and expe r imen ta l assistance. Thanks are also g iven to t h e Research M e m b e r C o m p a n i e s of The W e l d i n g Ins t i tu te w h o s u p p l i e d mate r ia l for e x a m i n a t i o n .

References

1. Castro, R., and de Cadenet, J. J., The Welding Metallurgy of Stainless and Heat Resisting Steels, Cambridge University Press, Cambridge, 1974.

2. Blom, U., and Helmer, C , "The Corro­sion Resistance of SANDVIK 3RE60 Versus

Austenitic Chromium-N icke l Steels and Ferritic Chromium Steels," SANDVIK Re­port 71-0201 33-11E, 1971.

3. Matejka, W. A., and Knoth, R. J., "A High Purity Approach to High Chromium Ferritic Stainless Steels," /. Testing and Eval­uation, 3 (3), May 1975, pp. 199-207.

4. Gooch, T. C , and Wi l l ingham, D. C , Weld Decay In Austenitic Stainless Steels, The Welding Institute, Abington, Cam­bridge, 1975.

5. Polgary, S., "Resistance to Intercrystal­line Corrosion in Stainless Steel Weld Metal , " Svetsaren (English Edit ion), 8 (1-2), 1972, pp. 2-12.

6. ASTM Standard A262-70, Modi f ied Practice E.

7. Gooch, T. G., "Corrosion of AISI Type 304 Austenitic Stainless Steel," Brit. Weld, j., 15 (7), July 1968, pp. 345-357.

8. Henrikson, S., and Knutsson, L., "Cor­rosion Tests In Baltic Seawater On Heat Exchanger Tubes of Various Metal l ic Mate­rials," Br. Corros. j., 10 (3), 1975, pp. 128-135.

9. Riggs, O. L., Jr., "Effects of Hydrogen Halides on Anodic Polarisation of Stainless Steel," Corrosion, 19, (5), May 1963, pp. 180t-185t.

10. Gooch, T. G., "We ld Decay: The Assessment of Susceptibil ity," Weld ing Institute Research Bulletin, 10 (11), Nos. 1969, pp. 287-292.

11. Rarey, C. R., and Aronson, A. H., "Pitt ing Corrosion of Sensitised Ferritic Stainless Steel," Corrosion, 28 (7), July 1972, pp. 255-258.

12. Hodges, R. I., "Intergranular Corro­sion In High Purity Ferritic Stainless Steels,"

Corrosion, 27 (3), March 1971, pp. 119-127 and Ibid., 27 (4), April 1971, pp. 164-167.

13. Bond, A. P., "Mechanisms of Inter­granular Corrosion in Ferritic Stainless Steels," Trans. Met. Soc, AIME 245, Oct. 1969, pp. 2127-2134.

14. Steigerwald, R. F., " N e w Molyb­denum Stainless Steels For Corrosion Resistance: A Review of Recent Devel­opments," Mat. Pert, 13 (9), Sept. 1974, pp. 9-16.

15. Wright, R. N., "Mechanical Behaviour and Weldabi l i ty of a High Chromium Ferritic Stainless Steel As a Function of Purity," Welding journal, 50 (10), Oct. 1971, Res. Suppl., pp. 434-s to 440-s.

16. Lowrie, R„ "Weld ing E-Brite 26-1 to Other Alloys," Weld ing Journal, 52 (11), Nov. 1973, Res. Suppl., 500-s to 506-s.

17. Di l lon, C P., "Use of Low Interstitial 26Cr-1Mo Stainless Steel in Chemical Plants," Mat. Pert., 14 (8), Aug. 1975, pp. 36-37.

18. Carlen, J, C , "Fabrication Plant in SANDVIK 3RE60," SANDVIK Report No. 71 0201 52-31 E.

19. Shaw, D., and Edwards, M., "A Tran­sistorised Potentiostat System for Corrosion Studies," Corrosion Science, 5 (6), June 1965.

20. Nordin, S., "Observations Regarding Stress Corrosion Phenomena in an Auste-no-Ferritic Stainless Steel," Uddeholm Steel Research Report, Achema-Congress, Frank-furt-am-Main, 1973.

21. Dorschu, K. E., "Weldabi l i ty of a New Ferritic Stainless Steel," Welding journal, 50 (9), Sept. 1971, Res. Suppl., pp. 410-s to 418-

WRC Bulletin 229 August 1977

(1) Dynamic Fracture-Resistance Testing and Methods for Structural Analysis

by E. A. Lange

The potent ia l fo r the in i t i a t ion of f as t f r ac tu re can be p red ic ted by the recent ly developed techno logy of l inear elast ic f r ac tu re mechan ics (LEFM) wh ich has p roduced the basis for an ana ly t ica l app roach to f rac tu re res is tance and s t ruc tu ra l in tegr i ty . To m a k e f r ac tu re mechan i cs a v iab le eng inee r i ng des ign too l , emp i r i ca l cor re la t ions between pract ica l d y n a m i c tes t resul ts a n d the basic pa rame te rs are needed. In th is paper the a t t r i bu tes and l im i ta t i ons of the Charpy , Drop Weight -N i l Duct i l i ty T rans i t i on T e m p e r a t u r e , Drop Weight Tear, a n d Dynamic Tear tes ts are d iscussed w i t h respect to p rov id i ng i n f o r m a t i o n usefu l in s t ruc tu ra l in tegr i ty analyses.

(2) Junction Stresses for a Conical Section Joining Cylinders of Different Diameter Subject to Internal Pressure

by W. J. Graff

The conica l t r ans i t i on sect ion of a cy l indr ica l pressure vessel was i n s t r u m e n t e d inside and ou ts ide w i th electr ic res is tance s t ra in gages, and f r o m the l ong i t ud ina l a n d c i r c u m f e r e n t i a l s t ra ins measured exper imenta l l y , the co r respond ing stresses were d e t e r m i n e d . The resul ts were c o m p a r e d w i th ca lcu la ted stresses f r o m the theory of shel ls. The exper imen ta l l y d e t e r m i n e d st resses exceeded code m e m b r a n e stresses in the i m m e d i a t e v ic in i ty of t he j unc t i ons .

Pub l ica t ion of th is paper was sponsored by t he We ld ing Research Counc i l .

The pr ice of WRC Bul le t in 2 2 9 is $ 7 . 5 0 per copy. Orders shou ld be sen t w i t h p a y m e n t to t he We ld ing Research Counc i l , Un i ted Eng ineer ing Center , 345 East 4 7 t h Street , New York , NY 10017 .

W E L D I N G R E S E A R C H S U P P L E M E N T I 353 -s