Resistance to Hydrogen Cracking

in Steel Weld Metals

Microstructure at low H2 levels and hardness at high H2 levels

are determining factors in cracking resistance

BY P. H. M. HART

ABSTRACT. In comparison to HAZ hydro­gen cracking, there is little reported work on the influence of composition and microstructure on hydrogen cracking in weld metal. This paper describes the results of work using the G-BOP test to examine the effect of these factors at different heat inputs and hydrogen levels, mainly using SMAW, on the resistance to hydrogen cracking in low dilution weld deposits.

For C-Mn, C-Mn-Mo, C-Mn-Ni and C-Mn-Ni-Mo deposits at around 10 ml H2/IOO g deposited weld metal, the resis­tance to cracking was primarily controlled by weld deposit hardness, with increasing hardness (over the range -v 200-330 HV) indicating a decreased resistance to cracking.

At lower levels of hydrogen (<5 ml / 100 g), resistance to cracking was con­trolled much more by microstructure than by hardness. Microstructural types associated with good notch toughness (e.g., predominantly acicular ferrite) pro­moted good resistance, while predomi­nantly ferrite with aligned M-A-C micro-structures had poorer resistance to crack­ing.

Introduction

Certain characteristics of hydrogen cracks, e.g., their often tight, convoluted morphology and their tendency to be delayed after welding for a few hours or

Based on a paper presented at the 65th A WS Annual Convention held in Dallas, Tex., April 9-13, 1984.

P. H. M. HART is the Deputy Head of the Materials Dept. at The Welding Institute, Cam­bridge, England.

even a few days, conspire against their easy detection by nondestructive testing. It is not unknown, therefore, for them to be undetected, and occasionally their presence coincides with an unfavorable joint toughness and service conditions initiating catastrophic failure of structures. Consequently, it is particularly important to be able to avoid hydrogen cracking.

The basic factors which, when present, simultaneously produce hydrogen cracks are well known (Ref. 1), and develop­ments in steel technology which have led to major reductions in plate carbon and carbon equivalent levels have allowed the use of lower preheat levels for avoid­ing HAZ hydrogen cracking in carbon manganese steels. While there have been major strides over the years in lowering the amount of hydrogen introduced from welding consumables into the weld, changes similar to those occurring in the plate have not occurred in the weld metal composition. Thus, while in the past the welding procedures used to avoid HAZ hydrogen cracking generally also protected the weld metal, the lower risk of HAZ cracking that modern lower carbon equivalent plates provide has resulted in the weld metal often being "exposed," dictating the required pre­heat level. In fact, in some instances, the requirement to preheat has prevented the attainment of the full economic advantage provided by the low carbon equivalent plates.

In comparison to the vast amounts of research published on compositional and microstructural factors influencing HAZ hydrogen cracking, the reported work on weld metal is sparse. It was therefore necessary, as part of the effort to safely yet economically fabricate steels, to increase the understanding of composi­

tional and microstructural factors which influence the resistance to hydrogen cracking in weid metal.

The aspects which are covered in the work described here are as follows:

C-Mn Consumables Shielded Metal Arc Welding 1. The effect of carbon content 2. The effect of coating iron powder content and basicity

Submerged Arc Welding 1. The effect of manganese content 2. The effect of flux type

Low Alloy Deposits (for Shielded Metal Arc) 1. The effect of Mo content 2. The effect of Ni content 3. Alternative compositions for 700 N/ mm2 (101.5 ksi) yield strength depos­its.

The influence of these factors on resis­tance to weld metal hydrogen cracking has been studied (generally at hydrogen levels of 2 and 10 ml/100 g and 0.8 and 3.0 k j /mm, or 20.3 and 76.2 kj/in.) in low dilution deposits, using the G-BOP test (Ref. 2) as described below.

Experimental

Steel for the G-BOP base blocks was 50 mm (2 in.) thick BS4360 grade 50D of the composition shown in Table 1. All the shielded metal arc consumables were special commercial productions, so as to provide constant base compositions while the intended factors were varied, and were of 4 mm (%•> in.) diameter. Consumables within a given series were produced by the same manufacturer

14-s | JANUARY 1986

with the exception of the series examin­ing the effect of nickel. Here the oppor­tunity was taken to compare the behav­ior of the 2.5% nickel level deposit pro­duced by two different manufacturers. For submerged arc welds, commercially available consumables were used, includ­ing a basic, a calcium silicate and a man­ganese silicate flux. The SMAW elec­trodes were dried to give weld hydrogen levels of 2, 5 and 10 ml/100 g deposited weld metal when assessed using the method in British Standard 639-1976. The submerged arc consumables were evalu­ated for weld hydrogen levels using the technique in Ref. 3.

To obtain a low dilution test weld bead, grooves were machined in the G-BOP test blocks as shown in Fig. 1. These were filled using the relevant con­sumable so that the test bead was diluted approximately'10% by base material. A test weld was prepared by clamping together two halves of the G-BOP plates with a G clamp tightened to a torque of 57 Nm and depositing automatically a weld run across the central gap of the plates. These were 75 mm (3 in.) long for SMAW deposits and 125 mm (4.9 in.) for submerged arc welds. Welding condi­tions were 180 A and 22 V with appropri­ate travel speed to give arc energy units of 0.8 or 3.0 k j /mm (20.3 or 76.2 kj/in.) for SMA welding. For the submerged arc welds, the conditions were 650 A, 30 V and 390 mm/min (15.4 in/min) (3 k j / mm, or 76.2 kj/in.). After approximately 20 hours, the region of the weld bead across the gap was heated with a gas torch to dull red to heat tint the surface of any crack formed from the root of the weld. Subsequently, the clamp was removed, and the test weld was frac­tured across the gap by simple bending. The fracture face of the weld was exam­ined at magnifications of up to 30X with the aid of a binocular microscope, and the proportions of tinted area (crack surface) were estimated visually. Dupli­cate test welds were usually carried out with block temperatures of 20°, 50°, 100° and 150°C (68°, 122°, 212°, and 302 °F) and for the high strength deposits up to 250°C (482°F), as necessary to prevent cracking.

To examine metallographicaliy the cracking produced, one weld of each electrode type and arc energy condition which was not heat tinted was sectioned longitudinally from the 75 mm (3 in.) test weld and prepared for metallographic examination, using a 2% nital etchant. Selected crack surfaces were examined in the scanning electron microscope.

Results

Chemical Analyses

The results of chemical analyses for the

100mm

Clamp Gamp

Test bead Buttered region

Section A -A

Fig. 1- G-BOP test

test weld beads, as seen in Table 1, were obtained by spectrographic analyses tak­en directly from the ground surface of the test welds. These show satisfactory constant base compositions within each series. It can also be seen that the two nominal 2.5% Ni deposits had similar compositions.

G-BOP Test Results

A typical set of results is shown in Fig. 2. In keeping with previous G-BOP work (Ref. 4), in order to characterize resis­tance to cracking of a particular weld, two variables were taken from the curve. These were the percentage cracking at the 20°C (68°F) preheat and the 10% crack preheat temperature (10% CPT). Generally, these gave the same ranking of behavior, but because at higher hydro­gen levels a number of deposits showed 100% cracking at a preheat of 20°C (68°F), the 10% CPT will mostly be used in discussing results. All the results are summarized in Tables 2-7.

Effect of Carbon Content

The results show a decreasing resis­

tance to cracking as the carbon content is increased, although the 10% CPT only rose by 25 °C (45 °F) as the carbon increased from 0.03 to 0.09%. The deposit hardness rose from 250 to 303 HV5 as the microstructure was changed by the refinement of the acicular ferrite and reduction of the amount of grain boundary ferrite. The crack morphology was typical of weld metal hydrogen cracking and showed no preference for the grain boundary ferrite rather than the transgranular locations —Fig. 3.

Iron Powder Series

Only small changes in resistance to cracking were noted as the percentage of iron powder or coating basicity was changed. There were small changes in deposit composition, principally in man­ganese, which were reflected in varia­tions in weld hardness. When the results are expressed versus weld metal hard­ness, as in Fig. 4, with other results at this hydrogen level and arc energy, it can be seen that there is no consistent effect from the iron powder additions and coat-

WELDING RESEARCH SUPPLEMENT 115-s

Table 1—Chemical Analysis of Plate and Welds

Element Weight % Si Mn Ni Cr Mo Cu

O N ppm ppm

Steel Plate BS 4360-50D

Test Weld Deposits

Mo and Ni Series

1.0 Mn 0.25 Mo 0.5 Mo 1.0 Mo 1.0 Ni 2.5 Ni 3.5 Ni 2.5 Ni "2nd"

Iron Powder Series

0.8 k j /mm

2.5 Bl, 0"„ Fe powder 2.5 Bl, 15% Fe powder 2.5 Bl, 30°b Fe powder 1.8 Bl, 25% Fe powder 3.0 Bl, 15% Fe powder

3.0 k j /mm

2.5 Bl, 0",7 Fe powder 2.5 Bl, 15% Fe powder 2.5 Bl, 30% Fe powder 1.8 Bl, 15% Fe powder 3.0 Bl, 15% Fe powder

Carbon Series

C1 C2 C3 C4

700 N/mm 2 Series

Ni-Cr-Mo 0.8 k j /mm 3.0 k j /mm

Ni-Mo 0.8 k j /mm 3.0 k)/mm

Cr-Mo 0.8 k j /mm 3.0 k j /mm

Submerged Arc Series

Basic Manganese silicate/S2 Calcium silicate/S2

nd —not determined.

0.17

0.06 0.07 0.07 0.07 0.06 0.05 0.06 0.05

0.07 0.05 0.05 0.05 0.06

0.06 0.05 0.06 0.04 0.04

0.03 0.05 0.06 0.09

0.07 0.06 0.06 0.06 0.08 0.07

0.10 0.05 0.04

0.029

0.008 0.009 0.009 0009 0.007 0.006 0.008 0.007

0.013 0.015 0.010 0.013 0.011

0.012 0.012 0.012 0.011 0.008

0.010 0.011

<0.005 0.006

0.011 0.010 0.012 0.009 0.010 0.009

0.013 0.015 0.022

0.022

0.012 0.012 0.013 0.012 0.012 0.010 0.012 0.015

0.017 0.017 0.014 0.015 0.018

0.018 0.015 0.016 0.015 0.015

0.012 0.012 0.009 0.010

0.020 0.021 0.022 0.015 0.019 0.018

0.017 0.036 0.013

0.40

0.31 0.45 0.38 0.37 0.34 0.34 0.39 0.42

0.39 0.35 0.34 0.33 0.32

0.31 0.26 0.32 0.29 0.24

0.28 0.31 0.31 0.32

0.45 0.38 0.45 0.40 0.46 0.33

0.23 0.49 0.43

1.30

1.06 1.15 1.01 1.10 1.08 1.04 1.09 1.11

1.61 1.55 1.65 1.59 1.73

1.48 1.40 1.68 1.53 1.49

1.60 1.73 1.67 1.73

1.62 1.58 1.07 1.01 1.22 1.09

0.96 1.49 0.76

0.05

0.04 0.05 0.05 0.06 0.90 1.83 2.81 1.94

0.04 0.04 0.03 0.04 0.05

0.04 0.04 0.04 0.04 0.03

0.05 0.05 0.04 0.04

1.76 1.96 2.59 2.90 0.02 0.02

0.02 0.02 0.02

0.07

0.02 0.04 0.03 0.05 0.04 0.03 0.04 0.05

0.03 0.04 0.03 0.03 0.04

0.03 0.03 0.03 0.03 0.03

0.04 0.04 0.03 0.04

0.42 0.48 0.05 0.05 1.78 1.82

0.02 <0.01 <0.01

0.03

<0.01 0.21 0.39 1.04 0.02 0.01 0.02 0.01

<0.01 0.01

<0.01 <0.01 <0.01

<0.01 <0.01 <0.01 <0.01 <0.01

<0.01 <0.01 <0.01 <0.01

0.29 0.32 0.43 0.48 0.95 0.97

<0.01 <0.01 <0.01

0.06

0.06 0.03 0.03 0.04 0.03 0.03 0.03 0.03

0.06 0.06 0.05 0.05 0.6

0.06 0.05 0.05 0.05 0.05

0.02 0.02 0.02 0.02

0.02 0.02 0.03 0.02 0.03 0.03

0.14 0.12 0.12

0.029

<0.005 0.009 0.007 0.012 0.008 0.005 0.009 0.007

<0.005 <0.005 <0.005 <0.005 <0.005

<0.005 <0.005 <0.005 <0.005 <0.005

<0.005 <0.005 <0.005 <0.005

0.006 0.003 0.007 0.006 0.008 0.006

<0.005 <0.005 <0.005

<0.01

<0.01 <0.01 <0.01 <0.01 <0.01

0.006 <0.01 <0.002

0.014 0.014 0.012 0.009 0.013

0.011 0.009 0.012 0.008 0.011

0.009 0.011 0.009 0.011

0.014 0.012 0.015 0.012 0.016 0.011

<0.010 <0.010 <0.010

0.007

<0.005 0.008 0.007 0.009 0.007 0.006 0.008 0.007

<0.005 <0.005 <0.005 <0.005 <0.005

<0.005 <0.005 <0.005 <0.005 <0.005

<0.005 <0.005 <0.005 <0.005

<0.003 <0.003 <0.003 <0.003 <0.003 <0.003

0.016 0.006 0.007

nd

nd nd nd nd nd 370 nd

490

345 395 327 416 325

316 328 264 345 268

354 343 335 328

320 240 330 250 300 250

nd nd nd

nd

nd nd nd nd nd 98

nd 65

122 122 77 77 58

167 77 55 68 66

84 156 167 162

130 130 110 120 130 110

nd nd nd

ing basicity. No unusual features were observed on the crack fracture surfaces or in their morphologies.

IIIII

so

a

r.0

20

I,

D

\

• \

\

ffl + " ~ - ^ ^ ^ O ^ * *

V "̂̂ * ateB

t 0 2 1 - Mn O rj 39 ,Mo

a 1J04%Mo

-- . a , ^

Preheat temperature ,"C

Fig. 2-Typical set of results: G-BOP test results. Mo series, 0.8 k//mm, 5 ml H2/100 g deposited metal

Effect of Molybdenum

At 2 ml H2/IOO g, no cracking was obtained, even at approximately 1.0% Mo. At the intermediate hydrogen level of 5 ml/100 g, cracking was obtained, but there were indications that at 0.8 and 3.0 k j /mm (20.3 and 76.2 kj/in.) there was an improvement in cracking resis­tance with 0.2 and 0.4% Mo additions when compared with results (Ref. 4) at 0% Mo. At 10 ml H2/100 g, increasing Mo was detrimental. A 1.0% Mo level raised the 10% CPT by about 50°C (90°F). These results are shown in Fig. 5.

Microstructurally, the most marked effect of increasing Mo was the well documented one of reducing the amount of grain boundary ferrite and ferrite with aligned M-A-C and promoting and refin­

ing acicular ferrite. This was most marked for the initial 0.2%-Fig. 7. At 1% Mo the boundary ferrite was virtually absent. The predominant crack morphology was transgranular, showing no preference for grain boundary ferrite.

Table 2 -Seriesw

Carbon content wt %

0.03 0.05 0.06 0.09

-G-BOP Test Results for Carbon

% Cracking at 20°C

(to nearest 5%)

30 30 95

100

10% CPT (to nearest

5QC)

75 80 95 95

Mean hardness

HV5

250 268 295 303

l , l0.8 k|/mm. 10 ml H2/ 100 g deposited metal

16-s | JANUARY 1986

Table 3—G-BOP Test Results for Fe Powder Series**'

% Cracking Electrode at 20°C(b> 10% CPT(c)

0.8 kj/mm

2.5BI 0% 100 90 Fe powder

2.5BI 15% 60 70 Fe powder

2.5BI 30% 100 90 Fe powder

1.8BI 15% 80 85 Fe powder

3.0BI 15% 100 85 Fe powder

3.0 kj/mm

2.5BI 0% 30 75 Fe powder

2.5BI 15% 25 60 Fe powder

2.5BI 30% 40 85 Fe powder

1.8BI 15% 35 50 Fe powder

3.0BI 15% 50 80 Fe powder

,al10 ml H2/10O g deposited metal <wTo nearest 5",,. <c)To nearest 5°C

Effect of Nickel

In contrast to the Mo series, some cracking was observed at 2 ml H2/IOO g, but only at the highest nickel content of 2.8% — Fig. 6. However, the electrode from the second manufacturer of the nominal 2.5% nickel level had a consider­ably higher resistance to cracking, giving 95% cracking at 20°C (68°F) or a 10% CPT of 85°C (185°F).

At 5 ml H2/IOO g and low arc energy (0.8 kJ/mm/20.3 kj/in.), the resistance to cracking decreased continuously as the nickel level was raised. In contrast, at 3.0 kj /mm (76.2 kj/in.), the low and interme­diate nickel additions, 0.8 and 1.8%, showed an improved resistance (by ref­erence to earlier results [Ref. 4] at 0% nickel), while the highest nickel level (2.8%) showed a decrease in cracking resistance. At 10 ml H2/100 g, increasing nickel produced a continuous decrease in resistance to cracking at both high and low arc energies. At 0.8 k j /mm (20.3 kj/in.), the 10% CPT rose by 70°C (126°F) for the 2.8% nickel level. The nominal 2.5% Ni electrode produced by

Powder

M e a n

hardness

HV5

275

260

271

263

251

228

209

218

209

Table 4 — G - B O P Test Results for M o Series

Electrode nominal % Cracking % Mo at 20°C<a> 10°,

0.8 k| /mm, 2 ml H2/100 g 0.25 0 0.5 0 1.0 0

0.8 k j /mm, 5 ml H 2 /100g 0 15 0.25 15 0.5 15 1.0 90

3.0 k l /mm, 5 ml H 2 /100g 0 35 0.25 5 0.5 10 1.0 30

0.8 k j /mm, 10 ml H2/100 g 0 30 0.25 100 0.5 100 1.0 100

3 k j /mm, 10 ml H2/100 g 0 45 0.25 50 0.5 50 1.0 90

( a ,To nearest 5"... ( b ,To nearest 5 'C .

CPT<b>

<20 <20 <20

40 30 30 75

40 <20

20 40

55 90 90

100

45 45 75 95

Mean hardness

HV5

262 277 291

237 262 277 291

198 218 230 242

237 264 280 292

198 218 230 242

Table 5—G-BOP Test Results for Ni Series

Electrode nominal

% Ni

0.8 k|/mm 1.0 2.5 2.5<c' 3.5

0.8 k j /mm, 0 1 2.5 3.5

3.0 k j /mm, 0 1 2.5 3.5

0.8 k j /mm 0 1.0 2.5 2.5<c>

3.5

3 k j /mm. 0 1 2.25 3.5

(a>To nearest 5 t b )To nearest 5

% Cracking at 20°C<a> 10°;

2 ml H 2 /100g 0 0

95 55

5 ml H 2 /100g 15 30 80

100

5 ml H 2 /100g 35

5 5

80

10 ml H 2 /100g 30 60

100 100 100

0 ml H 2 /100g 45 50 70 90

"C.

CPT<b>

<20 <20

85 45

40 45 50 75

40 <20 <20

50

55 70 95

110

125

45 50

75 110

(c>Electrode produced by second manufacturer

M e a n

hardness

HV5

270 288 300 344

237 270 288 314

198 219 239 256

237 273 287 295

314

198 219

239

256

216 Table 6—G-BOP Test Results for 700 N/mm2 Deposits

Table 7—G-BOP Test Results for Submerged Arc Series, 3 kj/mm

Electrode nominal % Cracking

at 20°C<a>

Ni-Cr-Mo 2 ml H2/100 g 0.8 k|/mm 100 3.0 k j /mm 95

10 ml H2/100 g 0.8 k j /mm 100 3.0 k j /mm 100

Ni-Mo 2 ml H 2 /100g 0.8 k j /mm 95 3.0 k j /mm 85

10 ml H 2 /100g 0.8 kj/mm 100 3.0 k j /mm 95

Cr-Mo 2 m l H 2 / 1 0 0 g 0.8 k j /mm 100 3.0 k l /mm 100

10 ml H 2 / 100g 0.8 k j /mm 100 3.0 k j /mm 100

( a )To nearest 5"„. < w To nearest 5 ° C

10% CPT<b>

145 130

230 225

100 50

140 140

190 220

250 250

M e a n

hardness

HV5

342 310

336 292

323 296

330 278

343 309

350 307

Electrode nomina l % Crack ing

at 20 D C ( a ) 10% CPT'b>

Basic f lux (7.3 ml H 2 / 1 0 0 g)

S2 w i r e 10 20

S3 w i r e 15 65 S4 w i r e 35 80

Mean hardness

HV5

225 256 273

Calc ium silicate f lux (4.9 ml H 2 / 1 0 0 g)

S2 w i r e 45 95

53 w i r e 30 120 54 w i r e 30 130

188 207

220

Manganese silicate f lux (4.7 m l H 2 / 1 0 0 g) S2 w i r e 40 90 209

S3 w i r e 35 115

S4 w i r e 65 125

t aTo nearest 5 V (b>To nearest 5 ' C

225 240

• VV ' 7'7 ' ' " 7 . * , ' .... a

' S ' j A i ^ - - v : ^ - 5 ^

Fig. 3 — Cracking in 0.03 % carbon series depos­it, 0.8 kj/mm, 10 ml H2/100 g (320X)

WELDING RESEARCH SUPPLEMENT 117-s

200

ISO

100

so-

+ o D X

• A

• • 3

Mn series Mo series Ni 2nd 2.5%Ni Iron powder series Carbon series Ni-Cr-Mo Ni-Mo Cr-Mo

•X"

Mn C-Mn-Mo C-Mn-Ni C-Mn-Nr-Mo

200 220 2i0 280

Mean deposit hardness , HV5

Fig. 4 — 10% CPT versus weld deposit hardness, all results

KO

120

~X. Cj

1 1 1 1

+ 2ml H2/100g O 5ml H2/100g • 10ml H2/100g

20

120 O 5ml H2/100g • 10ml H2/100g

0 0 2 OU 0-6 08 Wt, % Mo

10 12 li 0 02 (al Wt, % Mo lb)

Fig. 5- 10% CPT versus Mo content: A) 0.8 kl/mm, B) 3.0 k//mm

04 06 Wt, % Mo

10

the second manufacturer, not tested at 5 ml H2, gave a resistance at 10 ml similar to its companion deposit, in marked con­trast to the behavior at 2 ml H2/100 g.

Except at the highest nickel content and lowest arc energies, where the deposit contained some martensite, increasing nickel produced only a small change in the microstructural constitu­ents, viz., a slight refining of the acicular ferrite and a slight reduction in the amount of grain boundary ferrite. The nominal 2.5% nickel deposit from the second manufacturer had a contrasting microstructure which was predominantly ferrite with aligned second phase and is compared with its counterpart in Fig. 7. The most likely compositional differences responsible for this are the higher Ti and lower oxygen content of the deposit giving acicular ferrite.

While the crack morphology was pre­dominantly transgranular, at higher Ni

levels and deposit hardness (~ 270 HV), cracking at grain boundary ferrite was more noticeable.

700 N/mm2 (101.5 ksi) Deposits

The results for the three compositional routes evaluated, Ni-Cr-Mo, Ni-Mo and Cr-Mo, show that for all conditions the Cr-Mo deposits had the lowest resistance to cracking, while the Ni-Mo had the highest. While lowering the hydrogen level reduces the risk of cracking, increas­ing the heat input generally only had a small beneficial effect. At low heat inputs, all deposits contained substantial propor­tions of martensite. The most was in the Cr-Mo, and the least was in the Ni-Mo deposit. The microstructure of the Cr-Mo deposit was easily distinguishable from that of the other two by the virtual absence of any grain boundary nucleated product, and by the fairly uniform pres­

ence of the low temperature transforma­tion product described by Watkinson (Ref. 5) as "bainitic martensite." At the high heat input, the microstructures from all three deposits showed the formation of predominantly nonmartensitic trans­formation products, as evidenced by the reduction in hardness, and in the Ni-Mo deposit, the structure was showing the presence of some acicular ferrite.

Submerged Arc Welds

For all three fluxes, increasing the man­ganese content produced increases in deposit hardness and increases in 10% CPT temperatures. The major contrast, however, was between the results with the basic flux and those with the other two fluxes, since although the weld hydrogen level and deposit hardness were higher, the resistance to cracking, as evidenced by 10% CPT, was signifi­cantly greater for the basic flux. This appeared to be associated with the marked differences in microstructure between the deposits from the basic flux, which were composed predominantly of boundary ferrite and acicular ferrite, and the deposits from the other two fluxes, which consisted extensively of ferrite with aligned second phase — Fig. 8.

Discussion

The Effect of Microstructure and Hardness

In addition to the specific objectives of determining the effect of individual ele­ments and factors on the resistance to weld metal hydrogen cracking, a general objective was to improve the under­standing of the effect of microstructure and hardness on this type of cracking. Earlier work (Ref. 4), the results of which are also included in Tables 2-7, indicated that at the higher hydrogen level (10 ml/100 g), resistance to cracking was primarily related to the hardness of the weld deposit, while at a lower hydrogen level (5 ml/100 g), the microstructural characteristics of the deposit appeared to be more important than hardness in influ­encing crack resistance. During the more extensive evaluation of different compo­sitions reported here, the results have continued to provide further support for that view. As can be seen from Fig. 4, all the results (with the exception of the Cr-containing deposits) at the 10 ml H2/ 100 g level showed a good relationship between resistance to cracking and hard­ness, extending over a wide range (200-330 HV5). At the 5 ml H2 level, a similar trend is apparent, although the scatter has increased. The results for lower hydrogen levels (5 and particularly 2 ml H2) indicate that hardness becomes pro­gressively less useful for indicating the resistance to cracking.

18-s | JANUARY 1986

%0

120

WO

o

~ r

0

2nd manufacturer ®,0 2ml H2/100g

10ml H2/100g + 2ml H2/100g o 5ml H2/100g • 10ml H2/100g J i L

120

100-

o

CI

5ml H2/100g 10ml H2/100g

(a)

Fig 6-

0 0-5 VO VS 2 0 25 30 35 WT,% Ni

10% CPT versus Ni content: A) 0.8 k//mm, B) 3.0 k//mm

10 20 30 (b) Wt.% Ni

At the lower hydrogen levels, it was increasingly apparent that microstructure was much more important than hardness in controlling the resistance to cracking, and there were several clear cases where a change to a microstructural type, or alloy content, which would be expected to produce a better fracture toughness {e.g., a change from ferrite with aligned second phase to predominantly acicular ferrite or an increase in nickel content), led to an improvement in resistance to cracking. The first of these is where an alloy addition (up to 0.4% Mo or up to 1.8% Ni) lowered the 10% CPT, either by improving the microstructure, giving a reduced amount of grain boundary fer­rite and increased amount of acicular ferrite, as in the case of the Mo addition, or by the presence of Ni —Figs. 5 and 6. The second case is from the comparison of the behavior of the two nominal 2.5% Ni deposits. In this instance, it was found that differences in electrode coating for­mulation had produced markedly differ­ent resistances to cracking. These differ­ences were related to the different microstructures that had been produced, albeit at very similar levels of nickel and other elements. (The compositional fea­tures considered of most importance for the change of microstructure were the higher titanium and lower oxygen level in the deposit producing predominantly acicular ferrite.) The behavior of these two deposits at the low hydrogen level (2 ml/100 g) showed that the predominant­ly ferrite with aligned second phase struc­

ture had given a very significantly lower resistance to cracking than the predomi­nantly acicular ferrite structure. In con­trast, at the higher level of hydrogen (10 ml/100 g), the deposits, which were of similar hardness, yet still of different microstructure, were showing very simi­lar resistances to hydrogen cracking.

The third case concerned the sub­merged arc results. Here, the acicular

' < • •

ferrite microstructures produced with the basic flux had a much higher resistance to cracking than the ferrite with aligned second phase structures produced with the other two fluxes. This was despite a significantly higher hydrogen level and higher deposit hardness.

In addition to finding that for the higher hydrogen levels weld hardness was a good indicator of cracking risk, it was also

Fig. 7 —Nominal 2.5% Ni deposit microstruc­tures, 0.8 k//mm (320X): A) first manufacturer, B) second manufacturer

Fig. 8 — Submerged arc deposit microstructure, 3.0 k//mm (320X): A) basic flux, B) manganese silicate flux

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WELDING RESEARCH SUPPLEMENT 119-s

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found that weld hardness could usefully be related to weld composition, even simply expressed as the traditional car­bon equivalent for heat-affected zone (IIW carbon equivalent). An improved and more widely applicable relationship was obtained by regressing composition and weld cooling time against weld hard­ness: Deposit hardness (HV5 = 427.5 CEH

+ 202.8 where CEH = C -F 0.09 Mn -F 0.05 Ni + 0.08 Cr + 0.09 Mo - 0.016 Atsoo-soo where C, Mn, Ni, Cr and Mo are deposit element content in wt % and At(8oo-5oo) is the weld cooling time in seconds. The multiple correlation coefficient was 0.97. The limits of the variables studied were C 0.03/0.09, Mn 1.0/1.7, Ni 0/2.9, Cr 0 / 1.8, Mo 0/1.0, At 3/10.

Cracking Morphology and Fracture Surface

The metallographic and fractrographic examinations have provided data on the morphology and fracture surface fea­tures of known weld metal hydrogen cracking. It was found that small micro­cracks, located ahead of the main crack, could be situated in either grain boundary ferrite or intragranular microstructures, as shown in Fig. 9, and that, at least for hardnesses below approximately 270 HV5, the crack path was not preferential­ly located in grain boundary ferrite.

Above approximately 270 HV, there was an increasing tendency for cracking to be located preferentially at grain boundary ferrite, and when the grain boundary ferrite was absent, smooth intergranular cracking was increasingly present.

Fig. 10 —Hydrogen crack fracture surfaces: A) micro-void coalescence type (2000X), B) cleavage type (1000X), C) intergranular type

20-s | JANUARY 1986

Fig. 11 - Characteristic hydrogen crack fracture surfaces: A) "rounded" quasi-cleavage type (2000X), B) "rounded" void coalescence type (1000X)

The work has shown that hydrogen crack fracture surfaces can take a wide variety of forms. Crack fracture surfaces observed covered the range from micro-void coalescence through cleav­age, to smooth intergranular — Fig. 10. Intermediate forms such as Fig. 11 are, in the author's experience, less typical of other forms of cracking and more partic­ular to hydrogen cracking only.

Avoiding Cracking

It was not the objective of this pro­gram to devise conditions (preheat, hydrogen level, heat input, etc.) for avoiding weld metal hydrogen cracking in particular fabrications. However, the pre­heat conditions for avoidance of cracking in the G-BOP test, while unlikely to be directly relevant to multipass butt joint

welds, because of the different rate of diffusion of hydrogen from single and multipass welds, may be useful as a guide for 1 or 2 run fillet welds and root runs in butt joint welds.

Multiple linear regression analysis of the data for welds at 10 and 5 ml H2

produced the equation: 10% CPT = 188.4 CEW - 108.3 where CEW = C + 0.378 Mn + 0.145 Ni + 0.468 Cr + 0.299 Mo - 0.012 Atgoo-soo + 0.039 H2

where C, Mn, Ni, Cr, and Mo are weld metal contents in wt % and Atsoo-soo is the weld cooling time. H2 is the standard (BS 639) weld deposit hydrogen level in ml/100 g deposited weld metal.

The fit of the data is quite good (Fig. 12), with a multiple correlation coefficient of 0.96. The limits of the variables studied were as above for hardness, and for H2 5

Fig. 12 —Regression line for composition, cool­ing time and weld hydrogen level against 10% CPT

to 10 ml. Comparison of conditions to avoid weld metal cracking found in tests by Jones and Hart (Ref. 6), with values predicted by the above equation, showed reasonable agreement for a sin­gle pass fillet weld (-v80°C/176°F pre­heat predicted versus 75°C/167°F observed), but unsafe prediction for a multipass butt joint weld (^35°C/95°F predicted versus 75°C/167°F observed). This limited comparison would lend sup­port for the suggested restricted applica­tion of the regression equation men­tioned above.

Apart from the specific effects of com­position on cracking presented previous­ly and covered by the above regression equation for hydrogen levels greater than approximately 5 ml, the present work has indicated a mild beneficial effect of nickel additions of up to 1.8% (deposit compo­sition) for higher heat inputs at the low hydrogen level (<5 ml). Small additions of Mo (up to -vO.4%) to a 1% Mn base composition also gave a mildly beneficial effect at low hydrogen levels. However, despite the benefits that an acicular fer­rite microstructure conferred on the resistance to cracking at low hydrogen levels, significant alloying eventually led to a reduction in resistance to cracking. At both 10 and 2 ml H2 levels the addition of Cr appeared to be markedly detrimen­tal, perhaps indicating an inherent detri­mental effect of the element over and above its hardening effect.

Summary and Conclusions

Using the G-BOP test, the influence on weld deposit composition and micro-structure on weld metal hydrogen crack­ing was examined. This examination was primarily for SMA welding, but some submerged arc deposits were also tested. Work was carried out over a range of hydrogen levels (2-10 ml/100 g deposit­ed weld metal) and heat inputs (0.8 to 3 kJ/mm/20.3 to 76.2 kj/in.). From the investigation, two general conclusions

WELDING RESEARCH SUPPLEMENT j 21-s

concerning the resistance to weld metal hydrogen cracking can be drawn:

1. At the higher hydrogen level of 10 ml/100 g, the resistance to cracking in SMAW C-Mn, C-Mn-Mo, C-Mn-Ni and C-Mn-Ni-Mo deposits was primarily con­trolled by weld metal hardness. At this hydrogen level, hardness was a useful indicator of the risk of cracking, with increasing hardness indicating an increas­ing risk of cracking. This was found to be true for hardnesses in the range of ~ 200-330 HV5.

2. At lower weld hydrogen levels (<5 ml/100 g), resistance to cracking of SMAW C-Mn, C-Mn-Mo, C-Mn-Ni and C-Mn-Ni-Mo deposits and submerged arc C-Mn deposits was controlled principally by microstructural type rather than hard­ness.

Microstructures (e.g., predominantly acicular ferrite) and alloy content (e.g., presence of Ni) which in combination would be expected to give good weld metal fracture toughness produced improved resistance to cracking. Pre­dominantly ferrite with aligned second phase microstructures gave reduced resistance to cracking.

In addition to the above general con­clusions, the following conclusions, spe­cific to stages of the investigation, were derived:

3. Increasing carbon from 0.03 to 0.09% decreased resistance to cracking in a 1.5% Mn deposit at 10 ml H2 /100g and 0.8 k j /mm (20.3 kj/in.).

4. Electrode coating iron powder con­tents up to 30% and changes in the coating basicity from 1.8 to 3.0 had no major effect on the resistance to cracking in 1.5% Mn deposits at the 10 ml H2/100

g level and arc energies of 0.8 and 3 kj /mm (20.3 and 76.2 k)/in.).

5. For a 1% Mn base composition, increasing the Mo content up to 1.0% significantly reduced the resistance to cracking at 10 ml H2/100 g and arc energies of 0.8 and 3.0 k j /mm (20.3 and 76.2 kj/in.).

At the 5 ml H2/100 g level, increasing the Mo content up to -^0.4% appeared to result in a small improvement in the resistance to cracking for both arc ener­gies studied, but 1.0% Mo reduced the resistance to cracking.

No cracking was observed at the 2 ml/100 g hydrogen level.

6. Increasing the Ni content up to 2.8% produced a progressive and signifi­cant fall in the resistance to cracking at 10 ml H2/100 g.

Increasing the Ni content to 2.8% pro­duced a progressive fall in the resistance to cracking at 5 ml H2/100 g and the low arc energy (0.8 k)/mm/20.3 kj/in.) condi­tion. At 3.0 k j /mm (76.2 kj/in.), Ni up to 1.8% appeared to increase the resistance to cracking.

At the 2 ml/100 g hydrogen level, cracking was observed only at the high­est (2.8%) Ni level.

7. Significant differences in resistance to cracking between the three composi­tions Ni-Cr-Mo, Ni-Mo and Cr-Mo for high strength (-v-700 N/mm2/105.5 ksi) deposits were observed. The lowest resistance was in the Cr-Mo deposit, and the highest was in the Ni-Mo deposit. At 10 ml H2/100 g, the Cr-containing depos­its had much lower resistance to cracking for a given hardness than all the other deposits tested, including the Ni-Mo.

8. The preheat to prevent cracking in

the G-BOP test was related to weld composition, hydrogen level and cooling time by the equation Preheat temperature (°C) = 188.4 CEW - 108.3 where CEW = C -F 0.378 Mn -F 0.145 Ni + 0.468 Cr -F 0.299 Mo - 0.012 Atooo-soo + 0.039 H2

This equation may serve as a guide for preheat temperatures to avoid weld met­al cracking in 1 or 2 pass fillet welds and root runs in butt joint welds, but is not applicable to multipass welds.

A ckno wledgments

The author would like to thank all the staff at The Welding Institute who have helped with the program, in particular Mr. G. T. Hall and Mr. N. Tebbit for their careful experimental work. Grateful acknowledgment is made to members of the sponsor group supporting this project and their supply of experimental consum­ables.

References

1. Hart, P.H.M. April 1982. Hydrogen in steels —its significance for welding. Presented at Hydrogen in Steel Conference, Institution of Metallurgists, Bath, England.

2. Graville, B.A., and McParlane, M. 1974. Weld metal cracking. Met. Con. 6 (2):62-63.

3. IIW Doc 11-1002-83. 4. Hart, P.H.M. 1982. Hydrogen cracking in

ferritic steel welds. Third Int. Conf. Hydrogen and Material, Paris, France.

5. Watkinson, F. 1969. Hydrogen cracking in high strength weld metals. Welding lournal 48 (9):417-s to 424-s.

6. Jones, A.R., and Hart, P.H.M. October 1983. Improving procedure 'prediction to avoid hydrogen cracking when welding C:Mn steels. Welding Institute Members Report 225/ 1983.

WRC Bulletin 306 July 1985

PVRC Flanged Joint User Experience Survey By J. R. Payne

Over 180 completed questionnaires were received f rom a distribution list of 295 flanged joint users. The results of the survey provided guidance in establishing research needs and priorit ies for flanged joint improvement programs.

The survey and the publication of the report were sponsored by the Task Group on Gasket Testing of the Subcommit tee on Bolted Flanges of the Pressure Vessel Research Commit tee.

The price of WRC Bulletin 306 is $14.00 per copy, plus $5.00 for postage and handling. Orders should be sent with payment to the Welding Research Council, Rm. 1301, 345 E. 47th St., New York, NY 10017.

22-s | JANUARY 1986