ARTICLE IN PRESS

Available at www.sciencedirect.com

journal homepage: www.elsevier.com/locate/he

I N T E R N A T I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 8

0360-3199/$ - see frodoi:10.1016/j.ijhyde

�Corresponding auE-mail address:

Hydrogen induced cold cracking studies on armour gradehigh strength, quenched and tempered steel weldments

G. Magudeeswarana,�, V. Balasubramaniana, G. Madhusudhan Reddyb

aCentre for Materials Joining Research (CEMAJOR), Department of Manufacturing Engineering, Annamalai University,

Annamalai Nagar 608 002, Tamil Nadu, IndiabMetal Joining Section, Defence Metallurgical Research Laboratory (DMRL), Kanchanbagh (P.O.) Hyderabad 560 058 Andhra Pradesh, India

a r t i c l e i n f o

Article history:

Received 2 October 2007

Received in revised form

29 January 2008

Accepted 29 January 2008

Available online 12 March 2008

Keywords:

Shielded metal arc welding process

Flux cored arc welding process

Austenitic stainless steel

Low hydrogen ferritic steel

Hydrogen induced cracking

Diffusible hydrogen

Implant test

nt matter & 2008 Internane.2008.01.035

thor. Tel.: +91 4144 239734magudeeswaran@yahoo.c

a b s t r a c t

Quenched and tempered (Q&T) steels are prone to hydrogen induced cracking (HIC) in the

heat affected zone after welding. The use of austenitic stainless steel (ASS) consumables to

weld the above steel was the only available remedy because of higher solubility for

hydrogen in austenitic phase. The use of stainless steel consumables for a non-stainless

steel base metal is not economical. Hence, alternate consumables for welding Q&T steels

and their vulnerability to HIC need to be explored. Recent studies proved that low hydrogen

ferritic (LHF) steel consumables can be used to weld Q&T steels, which can give very low

hydrogen levels in the weld deposits. In this investigation an attempt has been made to

study the influence of welding consumables and welding processes on hydrogen induced

cold cracking of armour grade Q&T steel welds by implant testing. Shielded metal arc

welding (SMAW) and flux cored arc welding (FCAW) processes were used for making welds

using ASS and LHF welding consumables. ASS welds made using FCAW process offered a

higher resistance to HIC than all other welds considered in this investigation.

& 2008 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights

reserved.

1. Introduction

Quenched and tempered (Q&T) steels are used in military

applications due to high hardness, high strength to weight

ratio and excellent toughness. These grades of Q&T steels are

prone to hydrogen induced cracking (HIC) leading to poor

ballistic performance [1–3]. The three methods of controlling

HIC Q&T welds are (i) temperature control method,

(ii) isothermal transformation method and (iii) the use of

austenitic stainless steel (ASS) weld metal. The temperature

control method depends on holding the weld at an elevated

temperature, in particular above that at which hydrogen by

diffusion is accelerated. The isothermal transformation

method prevents HIC by controlling the cooling rate of the

heat affected zone (HAZ) so that it transforms to softer (non-

tional Association for Hy

; fax: +91 4144 238080/238om (G. Magudeeswaran).

martensite) structure. During welding of Q&T steels, for

various reasons it is not possible to use preheat temperature

greater than 150 1C, hence the temperature control method is

severely restricted and isothermal transformation method

cannot be used. The only alternative is to use welding

consumables which virtually prevents the introduction of

hydrogen in HAZ and which produces a weld metal insensi-

tive to hydrogen [4].

ASS welding consumables are being used for welding Q&T

steels, as they have higher solubility for hydrogen in

austenitic phase, to avoid HIC. the same consumable finds

application for the welding of high hardness Q&T steels to

meet the service requirements in the construction of combat

vehicles [5]. The use of stainless steel filler for welding a non-

stainless steel base metal (BM) must be avoided as ASS fillers

drogen Energy. Published by Elsevier Ltd. All rights reserved.

275.

ARTICLE IN PRESS

I N T E R N AT I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 81898

are much more expensive. In the recent years, the develop-

ments of low hydrogen ferritic (LHF) steel consumables

that contain no hygroscopic compounds are utilized for

welding of Q&T steels [6,7]. The above practice paved a

new way for cost effective consumable selection to meet

out the requirements to avoid HIC during welding of

armour grade Q&T steels. The majority of armour fabr-

ication is performed by fusion welding processes and they

demand for high quality welds. Shielded metal arc welding

(SMAW) and the flux cored arc welding (FCAW) processes

are widely used in fabrication of combat vehicle construction

[8]. HIC in Q&T steel welds are influenced by (i) level of

diffusible hydrogen present in welds, (ii) tensile stresses

that act on the welds and (iii) susceptible HAZ micro-

structure [9]. Both welding process and welding consu-

mables have significant effect on the above factors that

influence HIC. Hence, in this investigation, an attempt

has been made to study the effect of welding consumables

(ASS and LHF) and welding processes (SMAW and FCAW) on

hydrogen induced cold cracking of armour grade Q&T steel

weldments.

2. Experimental work

2.1. BM and welding consumables

The BM used in this investigation is a Q&T steel, closely

confirming to AISI 4340 specification. The microstructure of

the BM exhibits acicular martensite (Fig. 1). In this investiga-

tion, ASS and LHF welding consumables were used to make

the welds by using SMAW and FCAW process. The weld

fabricated using ASS consumable and SMAW process is

referred as SA weld; the weld fabricated using LHF consum-

able and SMAW process is referred as SF weld. Similarly, the

weld fabricated using ASS consumable and FCAW process is

referred as FA weld; the weld fabricated using LHF consum-

Fig. 1 – Microstructure of the base metal.

able and FCAW process is referred as FF weld. Vacuum

spectrometer (ARLModel: 3460) was used to study the weld

metal and BM chemistry. Sparks were ignited at various

locations on the BM and weld metals and their spectra were

analyzed for estimation of respective alloying elements.

ASME, Sec IIC (2006) and ASTM E8M-06 guidelines were

followed for evaluating the mechanical properties of weld

metals and BM. The tensile test of the BM and all weld metal

was carried out in a 100 kN, electro-mechanical controlled

universal testing machine (Make: FIE-Bluestar, India; Model:

UNITEK-94100). The specimen was loaded at the rate of

1.5 kN/min so that tensile specimen undergoes uniform

deformation. The specimen finally fails after necking and

the load versus displacement was recorded. The 0.2% offset

yield strength was derived from the diagram. The percentage

of elongation was also determined. The chemical composi-

tion and mechanical properties of the base metal and all weld

metals are presented in Tables 1 and 2, respectively. The

welding process parameters used to fabricate the joints are

given in Table 3.

2.2. Diffusible hydrogen measurements

The important methods for the determination of diffusible

hydrogen content in weld metal are (i) mercury method, (ii)

glycerin replacement method, (iii) silicone oil replacement

method and (iv) gas chromatography method. Numerical

relations are available relating the amount of diffusible

hydrogen content determined by the above mentioned

different methods. However, the mercury method gives the

most reliable and repeatable results [10]. The diffusible

hydrogen levels in the weld metal of the welding consum-

ables were experimentally determined by mercury method

as per the guidelines dictated in the literature [11–13].

The diffusible hydrogen content of the weld metal sample

was made to collect over mercury at room temperature

for a sufficient time (72 h). The amount of hydrogen thus

released was measured by volumetric method using

a diffusible hydrogen measuring meter that had an inbuilt

gas burette for collecting the diffused hydrogen. The baro-

metric pressure as well as the precise temperature was

recorded. The volume of diffusible hydrogen per 100 g of the

deposited weld metal was calculated from the following

expression:

DH ¼ ½VgðB�HÞ=760� � ½273=ð273þ TRÞ�

� ½100=ðM2 �M1Þ�, (1)

where DH is the volume of diffusible hydrogen in ml/100 g of

deposited weld metal at NTP (0 1C and 760 mm Hg), Vg the

volume of gas in burette in ml after 72 h, B the barometric

pressure in mmHg, TR the room temperature 30 1C when Vg is

measured, H the head of mercury in mm at which Vg is

measured, M1 the mass of the sample in g before deposit of

the weld metal and M2 the mass of the sample in gm after

removal from hydrogen meter.

Five trials were carried out for each consumable and the

measured diffusible hydrogen values are presented in Table 4

along with mean and standard deviation of the above

measurements. The diffusible hydrogen in weld metal is

ARTICLE IN PRESS

Table 1 – Chemical composition of (wt%) of base metal and weld metal

Type of material Notationa C Si Mn P S Cr Mo Ni Fe

Base metal (Closely confirming to AISI 4340

grade)

BM 0.315 0.239 0.53 0.018 0.009 1.29 0.451 1.54 Bal

Austenitic stainless steel (Closely

confirming to AWS E307)

SA 0.099 0.56 6.59 0.022 0.008 19.614 2.68 9.18 Bal

Low hydrogen ferritic steel (AWS E11018-M) SF 0.050 0.242 1.30 0.020 0.014 0.133 0.222 2.12 Bal

Austenitic stainless steel (Closely

confirming to AWS E307 T1-1)

FA 0.073 0.93 6.04 0.014 0.001 19.854 0.005 8.36 Bal

Low hydrogen ferritic steel (AWS E110T5-K4) FF 0.042 0.280 1.23 0.009 0.009 0.54 0.51 2.21 Bal

a BM, base metal; SA, shieled metal arc welded austenitic stainless steel weld; SF, shieled metal arc welded low hydrogen ferritic steel weld; FA,

flux cored arc welded austenitic stainless steel weld; FF, flux cored arc welded low hydrogen ferritic steel weld.

Table 2 – Mechanical properties of base metal and all weld metals

Type of weld 0.2% yieldstrength (MPa)

Ultimate tensilestrength (MPa)

Elongation(%)

Base metal (BM) 1200 1290 12.5

Shieled metal arc welded austenitic stainless steel weld (SA) 660 735 35

Shieled metal arc welded low hydrogen ferritic steel weld (SF) 720 800 22

Flux cored arc welded austenitic stainless steel weld (FA) 565 600 30

Flux cored arc welded low hydrogen ferritic steel weld (FF) 680 760 15

Table 3 – Welding conditions and parameters

Parameters Unit Shieled metal arcwelded austenitic

stainless steel weld(SA)

Shieled metal arcwelded low hydrogenferritic steel weld (SF)

Flux cored arcwelded austenitic

stainless steel weld(FA)

Flux cored arc weldedlow hydrogen ferritic

steel weld (FF)

Pre heat

temperature

1C 100 100 100 100

Electrode

baking

temperature

1C

for

3 h

300 300 – –

CO2 gas flow

rate

l/

min

– – – 12

Filler

diameter

mm 4 4 2.4 1.6

Current A 170 160 260 220

Voltage V 26 23 35 30

Heat input kJ/

mm

0.88 0.85 1.5 1.3

I N T E R N A T I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 8 1899

usually expressed milliliters per 100 grams (ml/100 g) of

deposited weld metal or parts per million (ppm).

2.3. Implant testing

The implant test was conducted using implant testing

machine (Make: ISHA, India; Model: 27105) as per the

International Institute of Welding (IIW) guidelines [14–17]

and with modifications in base plate dimensions as detailed

in the literature [18]. An implant test system (Fig. 2) was

employed to evaluate the susceptibility of the material to HIC.

In this test system, a helical threaded specimen (fabricated

from the BM) was embedded in a single bead weld in a 14 mm

thick base plate. A single pass weld was deposited so that a

portion of the notch section was located in the HAZ. The

sample was subjected to the desired stress under constant

load within 5 min post-welding. This arrangement enabled

the coarse grained HAZ of the specimen to experience the

load. The time required for the implant specimen to fail under

each stress was noted and a plot of load–time was obtained

from this test. Three specimens were tested at each stress

level and the average of time to failure was used for plotting

ARTICLE IN PRESS

Table 4 – Diffusible hydrogen levels

Weld type Diffusible hydrogen content (DH) (ml/100 g) Mean diffusiblehydrogen

content (DH)

Standarddeviation for

DH

measurementsTrial 1 Trial 2 Trial 3 Trial 4 Trial 5 (ml/

100 g)ppma

Shielded metal arc welded austenitic

stainless steel weld (SA)

2.79 2.85 2.86 2.82 2.78 2.82 3.13 0.035

Shieled metal arc welded low

hydrogen ferritic steel weld (SF)

2.89 3.05 3.09 2.99 2.98 3.00 3.33 0.076

Flux cored arc welded austenitic

stainless steel weld (FA)

2.81 2.79 2.83 2.80 2.72 2.79 3.09 0.041

Flux cored arc welded low hydrogen

ferritic steel weld (FF)

2.92 2.99 2.98 2.89 2.97 2.95 3.27 0.043

a 1 ppm ¼ 1:11 ml=100 g of weld metal deposited.

Fig. 2 – Schematic diagrams of implant test system: dimensions in millimeters. (a) Test block (base plate); (b) implant

specimen; (c) test configuration.

I N T E R N AT I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 81900

stress–time curve (Fig. 3). From this plot, lower critical stress

(LCS) below which no failure was obtained and is presented in

Table 5.

2.4. Microstructure and hardness

The implant specimen that did not fail after 72 h was

subjected to microstructural examination to reveal the

presence of microcracks. Few implant specimens were

interrupted after 1000 min of loading and were subjected to

microstructural examination to reveal the crack path. The

microstructure analysis of the weldments was carried out

using a light optical microscope (Make: MEIJI, Japan; Model:

ML7100). The specimens were etched with 2% nital reagent to

reveal the microstructure of the weld region of LHF weld,

BM and HAZ regions. Aquaregia reagent was used to reveal

ARTICLE IN PRESS

I N T E R N A T I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 8 1901

the microstructure of the ASS weld region. Vickers’s micro-

hardness testing machine (Make: Shimadzu, Japan; Model:

HMV-T1) was employed for measuring the hardness in the

Table 5 – Lower critical stress (LCS) values

Weld type Lower criticalstress (MPa)

Shieled metal arc welded austenitic

stainless steel weld (SA)

420

Shieled metal arc welded low hydrogen

ferritic steel weld (SF)

350

Flux cored arc welded Austenitic stainless

steel weld (FA)

470

Flux cored arc welded low hydrogen

ferritic steel weld (FF)

370

420 MPa

350 MPa

470 MPa

370 MPa

250

350

450

550

650

750

1 10 100 1000 10000Failure Time (min)

App

lied

Stre

ss (M

Pa)

SASFFAFF

Fig. 3 – Implant test results: horizontal arrows indicate

specimen did not fail.

Table 6 – Microhardness variation across the weld

Weld type/region

Weldregion

Region close to theinterface in the weld

Shieled metal arc welded

austenitic stainless steel weld

(SA)

261 355a

Shieled metal arc welded low

hydrogen ferritic steel weld (SF)

311 394b

Flux cored arc welded austenitic

stainless steel weld (FA)

245 320a

Flux cored arc welded low

hydrogen ferritic steel weld (FF)

294 380b

a Grain boundary (white phase) phase region.b Region of hard untempered martensite.

weld metal region, fusion boundary, the region close to the

fusion boundary on weld metal side and HAZ region. ASTM E

384-05a guidelines were followed for measuring the micro-

hardness and the values are presented in Table 6. The

fractured surface of the implant specimen was analyzed

using scanning electron microscope (Make: JEOL, Japan and

Model: 5610LV) at higher magnification to study the nature of

fracture.

3. Results

3.1. Diffusible hydrogen level

In this study, the diffusible hydrogen levels of the four

consumables were determined by mercury method and the

results are presented in Table 4. The SA weld metal had a

diffusible hydrogen level of 2.82 ml/100 g of metal deposited

and while the SF weld metal recorded 3.00 ml/100 g of metal

deposited. Similarly, the FA weld metal had diffusible hydro-

gen level of 2.79 ml/100 g of weld metal as against 2.95

ml/100 g of weld metal by FF weld. The diffusible hydrogen

level did not show larger variations in all the four weld metals

considered in this investigation. However, the ASS welds

(SA and FA) showed relatively a lower level of diffusible

hydrogen level than LHF welds (SF and FF).

3.2. Lower critical stress (LCS)

It has been reported that there were cases of presence of

microcracks in those specimens that did not fail under

implant conditions even after 72 h [18]. In this study the

implant specimen that did not fail after 72 h was subjected to

metallographic examination to reveal the presence of micro-

cracks. From the Fig. 4, it is inferred that no microcracks were

found in the interface of all the welds. Hence, LCS was taken

as the stress below which no microcracks were present and

Mean (Hv) 0.5 kg load

weld/HAZregion side

Weld/HAZ interfaceboundary (fusion

boundary)

HAZregion

Basemetal

420 435 456

434 439 454

405 425 455

415 430 455

ARTICLE IN PRESS

µµm50µm50

50µm 50µm

HAZHAZ

WeldWeld

WeldWeld

HAZ

HAZ

GBP

GBP

UTMUTM

Fig. 4 – Optical micrographs of fusion zone of unfailed

implant specimen revealing non-existence of microcracks.

GPB: grain boundary phase; UTM: untempered martensite.

(a) SA weld; (b) FA weld; (c) SF weld; (d) FF weld.

20µµm 20µm

20µm20µm

Fig. 5 – Optical micrographs HAZ region close to the

interface of weld/HAZ interface of unfailed implant

specimen. (a) SA weld; (b) FA weld; (c) SF weld; (d) FF weld.

50µµm

AusteniteAusteniteDelta Ferrite

Delta Ferrite

Acicular FerritePolygonal Ferrite

50µm

50µm50µm

Fig. 6 – Microstructures of weld metal region. (a) SA joint; (b)

FA joint; (c) SF joint; (d) FF joint.

I N T E R N AT I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 81902

also the highest stress at which the fracture did not

occur (after 72 h). From Table 5 and Fig. 3, it could be in-

ferred that the welds made by ASS consumable showed a

higher LCS values than the LHF welds, irrespective of the

process used.

3.3. Optical micrographs

The implant specimen that did not fail after 72 h loading was

subjected to metallographic examination. The metallographic

examination revealed the non-existence of microcracks

(Fig. 4). The fusion zone (weld-HAZ interface region) in SA

and FA welds reveal a grain boundary phase (GBP) (Figs. 4a

and b). The width of GBP is larger for FA welds compared to

that of SA welds. The fusion zone microstructure of SF and FF

welds consist of untempered martensite (Figs. 4c and d). In all

the joints, the HAZ region in close proximity to the interface

invariably consists of untempered martensite (Figs. 5a–d).

However, the degree of fineness is not similar in all cases. The

HAZ region close to interface of SF weld (Fig. 5c) contains very

fine untempered martensite than SA weld (Fig. 5a). Similarly,

HAZ region close to fusion boundary of FA weld (Fig. 5b)

shows coarser untempered martensite than that of FF welds

(Fig. 5d). It is also evident from Fig. 5 that the SA weld (Fig. 5a)

contains finer untempered martensite than that of FA weld

HAZ region close to fusion boundary (Fig. 5b). Similarly, the SF

weld (Fig. 5c) reveals finer untempered martensitic feature

than that of FF weld (Fig. 5d) in HAZ region close proximity to

fusion boundary.

The micrographs of weld metal regions of all the joints are

displayed in Fig. 6. The weld metal region of the ASS joint

exhibits a skeletal delta ferrite in plain austenitic matrix

(Figs. 6a and b). However, the morphology of the delta ferrite

in the SA and FA welds are not the same. The FA weld exhibits

much widely spaced delta ferrite in a plain austenitic matrix

(Fig. 6b) whereas the SA welds exhibit much closely

embedded delta ferrite in a plain austenitic matrix (Fig. 6a).

ARTICLE IN PRESS

I N T E R N A T I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 8 1903

The weld metal region of SF joint exhibits fully acicular ferrite

morphology (Fig. 6c) whereas FF weld metal shows polygonal

ferrite matrix (Fig. 6d).

The micrographs taken at the fusion zone of the inter-

rupted (at 1000 min) specimen are displayed in Fig. 7.

It is evident from the figure that the crack is developed

in the interface of the weld/HAZ region and it is directed

toward the HAZ region which is characterized by high

hardness untempered martensite invariably in all the

welds. In ASS welds, the crack is developed in the GBP

(white phase) region (Figs. 7a and b). The crack is pre-

dominately located in the region of coarse untempered

martensite in the weld/HAZ interface region of LHF welds

(Figs. 7c and d).

3.4. Hardness

The hardness across the weld cross-section was measured

using Vickers’s microhardness testing machine and the

values are presented in Table 6. The hardness of the

unwelded BM is 455 Hv. The SA weld exhibits a hardness of

261 Hv in the weld metal region, while FA weld recorded 245

Hv. Similarly, the weld metal hardness is found to be 311 VHN

and 294 Hv for SF and FF welds, respectively. Similarly, the

hardness in the region adjacent to the fusion boundary in the

weld metal side (cracked region) are 355, 394, 320 and 380 Hv

for SA, SF, FA and FF welds, respectively. The hardness of the

fusion boundary of SA weld is 420 Hv while the SF welds

exhibit a hardness value of 434 Hv. Similarly, FA and FF welds

exhibit a hardness value of 405 and 415 Hv, respectively, in the

fusion boundary region. The hardness in the HAZ region of

50µµm

HAZ HAZ

HAZ

HAZWELD

WELD

WELD

GPB GPB

CRACKCRACK

CRACK CRACK

UTMUTM

UTM

UTM

WELD

50µm

50µm50µm

Fig. 7 – Optical micrographs of fusion zone of interrupted

implant specimen revealing existence of cracks in the

fusion boundary. GPB: grain boundary phase; UTM:

untempered martensite. (a) SA weld; (b) FA weld; (c) SF weld;

(d) FF weld.

the SA, SF, FA and FF welds are found to be 435, 439, 425 and

430 Hv, respectively. Thus, the weld made using ASS consum-

ables have a lower hardness in the weld region, region close to

the fusion boundary (where cracks are found), fusion

boundary and the HAZ region than their LHF steel counter-

parts.

3.5. Weld metal strength

The mechanical properties of the BM and all weld metals

are presented in Table 2. It is inferred that the BM has

1200 MPa yield strength and 1290 MPa ultimate tensile

strength. The yield strength of the SA, SF, FA and FF weld

metals are found to be 660, 720, 565 and 680 MPa, respectively.

The ultimate tensile strength of the SA, SF, FA and FF weld

metals are found to be 735, 800, 600 and 760 MPa, respectively.

The percentage elongation of the BM, SA, SF, FA and FF are

found to be 12.5%, 35%, 22%, 30% and 15%, respectively. Thus,

the LHF steel welds exhibit higher strength than the ASS

welds.

3.6. Fractured surface

The fractured surface of the failed implant specimen is

displayed in Fig. 8. An intergranular fracture is featured in

SA and FA implant specimen (Figs. 8a and b). On the other

hand quasi-cleavage type of fracture is featured in the

fractured surface of the SF and FF implant specimen

indicating brittle fracture (Figs. 8c and d). The fracture

surface analysis indicate that a higher energy fracture has

occurred in ASS welds under implant conditions than their

LHF counterparts.

Fig. 8 – Fractographs of the implant specimen. (a) SA welds;

(b) FA welds; (c) SF welds; (d) FF welds.

ARTICLE IN PRESS

I N T E R N AT I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 81904

4. Discussion

In the present investigation, the results indicate that the

welds made using ASS consumables have higher LCS and

thus they offer better resistance to HIC than the welds made

using LHF consumables, irrespective of the processes used.

The ASS welds made using FCAW weld has a better resistance

to HIC compared to SMAW counterparts. Similarly, the LHF

welds fabricated using FCAW process shows has a greater

resistance to HIC than their SMAW counterparts. Thus, the

welding consumables and welding processes have a signifi-

cant effect on the HIC of armour grade Q&T steel joints and a

detailed discussion on this issue is presented in the following

sections.

4.1. Diffusible hydrogen levels

The hydrogen gets introduced into the weld metal during

fusion welding processes in many ways including moisture in

flux and atmosphere, organic substance in flux, hydrogen in

core wire steel, extra hydrogenous material such as moisture,

grease, organic compounds, paint, etc. [10]. Most of the

hydrogen that is trapped in weld metal is in supersaturated

solid solution in weld metal. Hydrogen dissolved in a steel

matrix is diffusible, there by causing hydrogen embrittle-

ment. One practical method of greatly reducing or eliminating

cracking caused by hydrogen is to use low-hydrogen type

electrodes. Such electrodes have mineral covering that is very

low in hydrogen producing constituents [19]. The type of

coating of the covered electrodes and flux cored wire is one of

the major contributing factors of lower diffusible hydrogen

level in the present study. The covered electrodes and the flux

cored wires used in this investigation are made up of basic

type of flux. Basic electrodes and cored wires have calcium

carbonate or other basic carbonate based covering. They also

contain fluorspar which encourages fluidity of the slag.

Because of its basic character, this type of covering in welding

consumables provides good protection, for all types of steels

and hence, a good density of the weld metal and good

mechanical properties. Because of the nature of products of

which they are made, basic coverings require storage, drying

and handling to ensure the lowest possible hydrogen content.

It is only subject to these conditions that the use of these

electrodes is recommended to avoid cold cracking in welding

of steels [20]. Hydrogen levels of 123 ml=100 g of the weld

metal are attainable for electrodes and cored wires after

carefully controlled storage, baking and shielding procedures

are followed [21].

The hydrogen content in weld metals from basic electrodes

is much lower, around 5 ml/100 g of weld metal. Rapid

development of low-hydrogen content basic electrodes has

taken place during the last few years and it is now possible to

have as low as 3 ml/100 g of weld metal in special cases. These

electrodes are usually packed in special, diffusion-tight boxes

so that the moisture content can be kept low for a very long

time. The hydrogen in the weld metal is mainly due to the

moisture of the coating. This has two sources. The first one is

crystalline water found in some minerals and the second

comes when silicates and other extrusion aids are added. The

coating is extruded onto the core wire and the coating paste

must therefore possess good rheological properties. The

moisture can be removed by drying at elevated temperatures

greater than 250 1C. The moisture content of about 0.15–0.30%

can be achieved if proper drying is carried out before welding

which in turn produce very low hydrogen concentrations in

the weld metal [22].

The cored wires are manufactured by forming a band into a

tube, filled with a powder mix, and drawing the tube to the

correct size. After the drawing operation, the band has

reduced in thickness and at the same time the diameter of

the wire has decreased due to compaction of the powder

inside the tube. The particles in the compacted powder stick

together and to the sides of the tube rather well, preventing

the powder mixture falling out of the tube. The filling in the

core can either be mainly iron powder, with alloying elements

and some deoxidants, giving a metal cored wire, or minerals

similar to those used for coating of covered electrodes, giving

a flux cored wire. There are two main types of flux cored

wires: rutile and basic. They behave much the same as their

equivalent covered electrodes. One important difference is

that no silicates are needed for binding the grains together in

the cored wires. Thus, the moisture content is much less and

cored wires usually give very low hydrogen. Hydrogen levels

less than 5 ml/100 g weld metal can be obtained from basic

cored wires [23].

The flux in the ASS and the LHF steel electrodes used in

this investigation has no hygroscopic compounds which

will absorb moisture when exposed to atmosphere. Further

more, the welds (SA and SF) made using the above electrodes

were carried out in a well-protected atmosphere. The

electrodes that were baked to 300 1C for 3 h and thus removing

all the moisture present in the electrodes. In case of flux

cored wires (FA and FF welds), the flux is present inside

the cored wire and its exposure to atmosphere is minimum

and the metal core acts as a protective cover to the flux.

Carbon dioxide gas was used as a shielding gas during

welding with LHF flux cored wire which also prevents the

weld pool from the moisture present in the atmosphere. The

protective cover for the flux from atmospheric exposure and

use of CO2 as shielding gas for making welds are the

significant reasons for lower diffusible hydrogen levels

in FF welds. However, the ASS flux cored wire is a self-

shielded wire that contains no hygroscopic compounds and

no shielding gas was used for making FA welds. The absence

of hygroscopic compounds and the basic flux used in the

consumables (electrodes and flux cored wires) are the main

reasons for the lower level of diffusible hydrogen levels in all

the welds.

The amount of diffusible hydrogen level is not only lower

for the ASS welds but also for the LHF steel welds. The

diffusible hydrogen level in all the weld metals are found to be

within the acceptable criteria. For armour grade Q&T steels,

the weld metal should have a low hydrogen content in the

weld, i.e. o4 ml=100 g of deposited metal [10] irrespective of

the consumable used. As per the American Society of

Mechanical Engineer’s boiler pressure vessel codes [24,25],

the maximum permissible diffusible hydrogen level in low

alloy ferritic LHF steel electrodes (AWS E11018-M) and flux

cored (AWS E110T5-K4) wires is 4 ml/100 g of deposited metal

ARTICLE IN PRESS

I N T E R N A T I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 8 1905

and it corresponds to H4 (lower level) optimal supplemental

diffusible hydrogen level designator. Thus, the diffusible

hydrogen levels for all the consumables considered in this

investigation are within the above said acceptable criteria.

Thus, it is evident from the above discussion that low

hydrogen ferritic flux cored consumables can also be used

for welding armour grade Q&T steels as they have very low

level of diffusible hydrogen in weld metal.

There is no appreciable variation (nearly constant) in the

level of diffusible hydrogen in all the four welds considered in

this investigation. The consumables were selected specifically

so as to attain a lowest possible diffusible hydrogen level in

the weld metals. In Q&T steels, where the problem of HIC is

extremely significant, cracking susceptibility has been corre-

lated both with material hardness and strength and with

specific microstructures in various regions of the welds. High

strength welds are more susceptible to HIC than low strength

welds. Steels that transform martensitically are particularly

susceptible, especially the higher-carbon alloys with twinned

martensitic structure [19]. Thus, it is evident from the above

facts that weld metal chemistry, microstructure, hardness

and weld metal strength have a significant effect on HIC of

armour grade Q&T steel welds and are explained in detail in

the following sections.

4.2. Effect of weld metal chemistry, microstructure andweld metal strength on HIC

The weld metal chemistry plays a major role in the formation

microstructure of the welds and hence it has a direct

influence on the HIC. Weld microstructures has a significant

effect on the HIC and cannot be ignored. It is always pre-

ferred to avoid martensitic structure in any form that is

susceptible to hydrogen induced cold cracking leading to a

catastrophic failure during welding. Carbon is recognized as

the alloying element with the greatest influence on the

hardenability and cold cracking susceptibility of steels,

but other alloying elements such as molybdenum and

chromium also have strong influences on hardenability.

While nickel also increases the hardenability of steels,

it has been shown to reduce the cold cracking susceptibility

[26]. Nickel in weld metal plays an important role in

microstructural control. However, there is no general

agreement regarding the amount and combination proportion

ratios of Ni and other alloying elements in weld metal [27].

The higher nickel content improves the toughness in two

ways: nickel reduces the ferrite content of the weld metal

(magnetic microsturctural phase and more brittle than

austenite), and the nickel addition increases the toughness

in fully austenitic compositions. A secondary benefit is

that nickel stabilizes austenitic structure against the forma-

tion of martensite (another magnetic microstructural phase)

[28]. In the present investigation, SA weld metal has 9.18

wt% of nickel, while FA weld metal has 8.36 wt% of

nickel. Whereas the SF weld has 2.12 wt% and FF weld has

2.21 wt% of nickel in the FF weld metal. Higher nickel

content promotes the formation of greater proportion of

austenitic phase in ASS and hence the solubility of diffusible

hydrogen in SA and FA welds are much higher than SF and FF

welds.

In the present context, the weld metal microstructure

is a significant factor that governs cold cracking susceptibility

of armour grade Q&T steel joints. The hydrogen diffusion

coefficient of austenitic phase in weld metal is much

lower than the ferritic phase. But the solubility of hydrogen

in austenitic phase is much higher than the ferritic phase.

Thus, an austenitic phase in weld metal can store high

hydrogen content that cannot move fast enough to the

fusion boundary due to slow diffusion rate of hydrogen in

austenite phase. The diffusion coefficient of monatomic

hydrogen at room temperature is approximately five

orders of magnitude lower in austenite matrix than in ferrite

matrix. Thus, diffusion of hydrogen is more rapid in

ferrite than austenite, but the solubility of hydrogen

is approximately 30 times higher than ferrite at room

temperature [29]. Thus, austenitic phase acts as barrier to

hydrogen escape. The ductility of austenite is also high,

thereby providing an ability to overcome restraints. Moreover,

austenite has an ability to tolerate heavy dilutions without

risk of forming a martensitic structure that is susceptible to

weld cold cracking [30].

The FA welds exhibit much widely spaced delta ferrite in a

plain austenitic matrix (Fig. 6b) in the weld metal region

whereas the SA welds exhibit much closely embedded delta

ferrite in a plain austenitic matrix (Fig. 6) in the weld metal

region. The delta ferrite provides a preferred path for crack

propagation in the presence of hydrogen and thus the closely

embedded delta ferrite in SA welds (Fig. 6a) has a lower

resistance to HIC than FA welds that has a widely spaced

(Fig. 6b) delta ferrite in a plain austenitic matrix. The larger

austenitic phase and widely spaced delta ferrite morphology

is one of the contributing factor for higher resistance to HIC of

FA welds than SA welds. The SF weld exhibits fully acicular

ferrite morphology (Fig. 6c) whereas weld metal of FF weld

shows polygonal ferrite matrix (Fig. 6d). As already stated, the

solubility of hydrogen in austenite is high and the diffusivity

of hydrogen in austenite is low. The greater solubility of

hydrogen in ASS weld metals offers an advantage over ferritic

weld metals in reducing the occurrence of cold cracking due

to hydrogen.

HIC is not sensitive to composition of BM and weld metal,

but to the strength of the BM and weld metal. The strength

will also affect the inherent resistance of the weld metal to

the cracking effects due to hydrogen. The resistance to HIC is

inversely related to strength: lower strength metals being

more resistant to HIC [31]. In general, higher the strength of

the weld, lower is the resistance to weld cold cracking [19,32].

From Table 6, it is revealed that the FA weld has a lower

hardness in the weld metal region than other welds and

hence it offers higher resistance to HIC. Also, it is evident

from Table 2, that the ASS welds (SA and FA) have lower yield

strength and tensile strength than LHF welds (SF and FF).

Thus, the higher yield strength of the weld metals is also one

of the major influencing factor for lowering the resistance of

LHF welds (SF and FF) than their ASS counterparts (SA and

FA). However, FA offers a much greater resistance to HIC

owing to the lower yield strength than other welds. Thus,

lower hardness and lower yield strength of weld metal are the

contributing factors for enhancing resistance to HIC in FA

welds than other welds.

ARTICLE IN PRESS

I N T E R N AT I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 81906

4.3. Effect of fusion zone characteristics on HIC

The HAZ adjacent to the weld is raised to a high temperature

during welding and subsequent rapid cooling (quenching) by

the surrounding parent metal causes hardening. Close to the

fusion boundary, the HAZ is raised to a sufficiently high

temperature to produce a coarse grain region. This high

temperature region, because of its coarse grain, is not

only more hardenable but also less ductile than the re-

gions further away from fusion boundary. It is the region in

which the greatest risk of cracking exists. As a general rule,

for Q&T steels, the harder the microstructure the greater is

the risk of cracking. Soft structure can tolerate more

hydrogen than hard microstructure before cracking

occurs [20].

In case of similar welds the microstructures of the weld

metals and HAZ will be unique and essentially the zone

adjacent to the fusion boundary is a characteristic feature of

the BM and weld metal. But this is not true in the case of

dissimilar welds. The fusion boundary microstructure in

dissimilar welds often possesses some unique features.

Normal epitaxial nucleation during solidification along the

fusion boundary gives rise to grain boundaries that are

continuous from the BM into weld metal across the fusion

boundary. These boundaries are roughly perpendicular to the

fusion boundary and have been referred to as ‘‘Type I’’

boundaries. In dissimilar welds, where an austenitic weld

metal and ferritic BM exist, a second type of boundary that

runs roughly parallel to the fusion boundary is often

observed. This has been referred to as a ‘‘Type II’’ boundary

[33]. These boundaries typically have no continuity across the

fusion boundary to grain boundaries in the BM. Several

investigators have reported that hydrogen-induced disbond-

ing typically follows Type II grain boundaries [34–36].

The fusion zone microstructure of the ASS welds (SA and

FA) exhibits a soft GBP (white phase) of the interface similar to

that of the Type II boundaries as described above. The region

of white phase and fusion boundary are located in close

proximity distance to each other. It is observed that the

micrographs of the interrupted implant specimen revealed

that the crack was developed in the GBP (white phase) region

(Figs. 7a and b) in the ASS welds. The formation of the white

phase (rich in carbon and chromium) in the FA and SA welds

is due to the diffusion of carbon from BM region to weld metal

region and migration of n chromium from weld metal region

to BM region [35] and it depends upon the weld thermal cycle

employed for fabricating the joints. In this study, the rate of

diffusion of elements (carbon and chromium) across the

fusion zone is higher for FA joint due to relatively higher heat

input (1.5 kJ/mm) compared to SA joint that recorded a lower

heat input (0.88 kJ/mm). The width of the GBP in FA weld is

relatively larger compared to SA welds (Figs. 4a and b and 7a,

and b). The microhardness values (Table 6) reveal that FA weld

has a softer GBP region compared to SA joints. The above

variations in the GBP features are due to the difference in the

heat input employed for fabricating the welds and the rate of

diffusion of elements. The fusion boundary hardness (Table 6)

is higher for SA joint compared to the fusion boundary

hardness of the FA joint due to the minor variations in the

heat input employed.

On the other hand, the fusion zone microstructure of the

LHF joints (SF and FF) has hard untempered martensite and

no Type II boundary exists adjacent to the fusion boundary.

The crack followed this region of untempered martensite

(Figs. 7c and d) in LHF welds. The formation of untempered

martensite is due to the diffusion of carbon from the BM to

the weld metal region and is greatly influenced by the heat

input employed for fabricating the LHF joints. FF welds

recorded relatively a higher heat input (1.3 kJ/mm) than the

SA welds (0.85 kJ/mm). Thus, the rate of diffusion of carbon

from the BM region to the weld metal region in the FF welds is

relatively higher compared to SF welds. This resulted in minor

variations in fusion zone characteristics (region of untem-

pered martensite and fusion boundary) and it is clearly

evident from the microhardness values in the above region

(Table 6). Thus, the hardness in the region of untempered

martensite and the fusion boundary is higher for SF joint

compared to FF joint.

The results of this study have shown that hydrogen

introduced during welding can lead to HIC in dissimilar joints

(ASS) and also similar joint (LHF). Hydrogen in the welding arc

is detrimental in two ways: (1) it increases dilution by the

carbon steel BM, increasing the amount of untempered

martensite formed and (2) it interacts with untempered

martensite under stress to cause cracking. The most suscep-

tible microstructure (untempered martensite) forms in the

vicinity of the fusion boundary [36], the cracks are found in

the regions near to the fusion boundary invariably in all cases

and also the implant specimens failed in the fusion boundary

invariably in all cases.

The hardness and microstructure in fusion zone (i.e. GBP in

the case of ASS welds and region of untempered martensite

in the case of LHF welds) are the most influencing factors that

contributes for greater LCS values and higher resistance to

HIC. The hardness in the fusion zone is inversely proportional

to the critical stress (LCS). The lower the hardness in the

fusion zone the higher is the LCS and greater will be

resistance to HIC. Thus, lower fusion zone hardness of the

ASS joints is one of the major contributing factors for higher

LCS values compared to their respective LHF joints. The FCAW

joints exhibited higher LCS values than their respective

SMAW joints and hence higher resistance to HIC. In the

present investigation, FA joint has a greater LCS values

(470 MPa) and it offers greater resistance to HIC than all other

joints.

4.4. Effect of HAZ microstructure on HIC

Several factors affect the susceptibility of the material to weld

cold cracking such as strength, microstructure and alloy

composition. It is difficult to separate the effects individually

because the three factors are interrelated [32]. The nature of

the fusion zone and the HAZ has a significant effect on the

HIC susceptibility of armour grade Q&T steel. Microstructure

(in the HAZ adjacent to the fusion boundary) is probably the

most important variable in controlling the susceptibility of

weld cold cracking. The susceptibility of steels to weld cold

cracking increases with strength of the weld and is also

usually associated with hard microstructures. However, at the

same hardness, different microstructures can have different

ARTICLE IN PRESS

I N T E R N A T I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 8 1907

susceptibilities for weld cold cracking. For example, low

carbon martensite/bainite structures have greater suscept-

ibility for cold cracking than fine grained acicular ferrite at

equivalent hardness levels. Further, it is generally accepted

that untempered, twinned or dislocated lath martensite are

the most susceptible microstructures, while the preferred

substructure combines a well-tempered martensite or bainite

with an ausworking process to produce a refined packet size

and a uniform dispersion of carbides [37]. It is evident from

the micrographs of HAZ region in close proximity (Fig. 5) to

the fusion boundary consists of invariably untempered

martensite in all cases. However, due to the smaller differ-

ence in the heat input of the welding process, the hardness in

this region show minor variations. The hardness in the

HAZ region of the SA weld metal is relatively lower than

SF. Also, the FA weld HAZ region has relatively lower hardness

than FF weld. However, the HAZ hardness of SA weld is much

higher compared to that of FA weld. Similarly, the HAZ

hardness of SF is greater than the FF weld. In general, hard

HAZ microstructure is more susceptible to HIC than soft

(coarse) microstructures. FA weld shows relatively lower

hardness in the HAZ region compared to all other welds.

Thus, lower HAZ hardness is also one of the contributing

factors for improved resistance of FA welds against HIC than

other welds.

4.5. Effect of residual stress

Owing to localized heating during the welding process and

subsequent rapid cooling, residual stresses can arise in the

weld and BM. Such stresses are usually of yield point

magnitude. Residual stresses attributed to welding pose

significant problems in the accurate fabrication of structures

because those stresses heavily induce brittle fracturing

and degrade the buckling strength of welded structures.

These residual stresses are developed in the vicinity of a

weld during arc welding. A weldment is locally heated by

most welding processes, therefore, the temperature distribu-

tion in the weldment is not uniform, and metallurgical

changes take place as welding progresses along the weld.

Therefore, the welding residual stress is sometimes called

restraint stress [38,39]. These stresses can give rise to

distortion and under certain circumstances even to prema-

ture failure.

Hydrogen-assisted cracking is often detected at a notch in a

weld made under restraint. Thus, residual stresses play an

important role as far as the quality and reliability of a welded

construction are concerned [40,41]. Lee et al. [42,43] in their

investigations revealed the occurrence of cold crack, referred

as transverse crack, was caused by a complex interaction of

the diffusible hydrogen supply, susceptible microstructure

and tensile residual stress. Recent investigations by Madhu-

sudhan Reddy et al. [44,45] revealed that the magnitude of the

residual stress was found to vary with respect to welding

processes. Their findings suggested that FCAW process was

associated with lower residual stress as compared to SMAW

process.

The effect of residual stresses on HIC cannot be ignored and

it has a greater significance. However, in this investigation

residual stresses on the welds were not measured but, their

significance cannot be ignored and it can be explained in

general terms of welding. In the present investigation

the heat input of the SA, FA, SF and FF welds are 0.88, 1.5,

0.85 and 1.3 kJ/mm, respectively (Table 3). There is an

appreciable variation in heat input of SA and FA welds and

the same trend is also observed with SF and FF welds.

Thus, the FCAW welds have a higher heat input than their

respective SMAW welds. The magnitude of the residual

stress is inversely proportional to heat input. The higher

the heat input the lower is the magnitude of residual

stress. Thus, FCAW welds will have lower residual stress

than SMAW welds. The most significance of the residual

stress with respect to HIC is that it has a directly proportional

relationship. Thus, higher the residual stress, the greater is

the risk of HIC. Hence, welds with lower magnitude of

residual stresses will be preferred to resist HIC. Hence, lower

residual stress due to higher heat input in FA welds is also one

of the contributing factors for greater resistance to HIC than

other welds.

In summary, much widely spaced delta ferrite in a large

plain austenitic matrix, lower diffusible hydrogen level, lower

weld metal strength, softer GBP, lower fusion boundary and

HAZ hardness and lower residual stress are the possible

reasons for the greater resistance of FA welds against HIC

compared other joints.

5. Conclusions

In this paper, the effect welding consumables on hydro-

gen induced cold cracking of armour grade Q&T steel welds

made by SMAW and FCAW processes was analyzed in detail.

From this investigation, the following conclusions are

derived.

1.

The welding consumables and welding processes have

significant effect on HIC of armour grade Q&T steel welds.

2.

The welds made using ASS consumables (SA and FA)

offered a greater resistance to HIC than their respective

welds made using LHF consumables (SF and FF).

3.

The welds made by FCAW process (FA and FF) exhibited

greater resistance to HIC than their respective SMAW

welds (SA and FF).

4.

The joint fabricated by FCAW process using ASS consum-

able exhibited superior resistance to HIC compared to all

other joints.

Acknowledgments

The authors are thankful to Armament Research Board

(ARMREB), New Delhi for funding this project work (Project

no. MAA/03/41), M/s Combat Vehicle Research Development

Establishment (CVRDE), Avadi, Chennai for providing base

material and Department of Manufacturing Engineering,

Annamalai University for providing testing facility and M/s

Defence Metallurgical Research Laboratory (DMRL), Hydera-

bad for providing the facility to carry out metallurgical

characterization for this investigation.

ARTICLE IN PRESS

I N T E R N AT I O N A L J O U R N A L O F H Y D R O G E N E N E R G Y 3 3 ( 2 0 0 8 ) 1 8 9 7 – 1 9 0 81908

R E F E R E N C E S

[1] Ade F. Ballistic qualification of armour steel weldments.Weld J 1991;70:53–4.

[2] Coe FR. Welding steels without hydrogen cracking. TheWelding Institute Report. Cambridge, England, UK: AbingtonPublishing; 1973.

[3] Alkemade SJ. The weld cracking susceptibility of highhardness armour steel. Australia: Defense Science andTechnology Organization; 1996. p. 1–14 AR no. 009-659.

[4] Coe FR. Hydrogen in weld metal. Cambridge, England, UK:Abington Publishing; 1969.

[5] Madhusudhan Reddy G, Mohandas T, Tagore GRN. Weld-ability studies on high-strength low-alloy steel using auste-nitic stainless steel filler. J Mater Process Technol1995;49:213–28.

[6] Deb P Challenger, Clark DR. Transmission electron micro-scopy characterizations of preheated and non-preheatedshielded metal arc weldments of HY-80 steel. Mater Sci Eng1988;77:155–67.

[7] Cary HB. Modern welding technology. London: Prentice-HallInc.; 1999.

[8] Wilson J. The development of welding techniques for britishfighting vehicles, UK Crown Copyright; 1978. p. 47–55.

[9] Stout RD. Weldability of steels. New York: Welding ResearchCouncil; 1987.

[10] Basu B. Special consideration for use of high strength steelweldments. In: Shah AK, editor. National conference on‘‘welding—productivity and quality,’’ organized by NavalMaterials Laboratory, Ambernath, India. Mumbai, India:Allied Publishers; 2007. p. 73–8.

[11] Standard methods for determination of diffusible-hydrogencontent of martensitic, bainatic and ferritic steels producedby arc welding. In: American welding society guidelines.AWS: 4.3; 1986.

[12] Method for determination of diffusible-hydrogen content ofdeposited weld metal from covered electrodes in weldingmild and low alloy steels. IS: 11802; 1986.

[13] Fabling FT, Chew B. Diffusible hydrogen analysis. Met ConstrBr Weld J 1972:64–7.

[14] Recommendations for the use of the implant test as onadditional information on cold cracking susceptibility whenwelding steels (IIS/IIW-447-73). Weld World 1974;1(2):9–16.

[15] Sawhill JM, Dix AM, Savage WF. Modified implant test forstudying delayed cracking. Weld J 1974;53(12):554s–60s.

[16] Granjon H. Met Constr. Br Weld J 1969;1(11):505–15.[17] Destructive tests on welds in metallic materials—cold

cracking tests for weldments—arc welding processes—part3: externally loaded tests. ISO 17642-3, 2005.

[18] Ramasubbu V, Albert SK, Guruprasad V, Bhaduri AK. Evalua-tion of the hydrogen assisted cracking (HAC) susceptibility ofmodified 9Cr-1Mo (P91) steel using implant test. In: Nationalwelding seminar (NWS 2006), organized by Indian Institute ofWelding, Chennai, India, 2006. p. 342–9.

[19] Steve L. Weld integrity and performance. USA: ASM Interna-tional; 1997.

[20] Granjon H. Fundamentals of welding metallurgy. Bombay:Jaico Publishing House; 1994.

[21] Dobly RE. Joining issues for new constructional steels. SciTechnol Weld Joining 2000;5(6):341–6.

[22] Larks-Erik S. Control of microstrucres and properties in steelarc welds. Boca Raton: CRC Press; 1994.

[23] Norman B. Weldability of ferritic steels. Cambridge, UK:Ambington Publishing; 1994.

[24] Specifications for low alloy steel electrodes for shieldedmetal arc welding. American society of Mechanical Engi-neer’s boiler pressure vessel code. II C-SFA 5.5; 2007. p. 105–7.

[25] Specifications for low alloy steel electrodes for flux coredmetal arc welding, American society of Mechanical Engi-neer’s boiler pressure vessel code. II C-SFA 5.29; 2007.p. 641-86.

[26] Rowe MD, Nelson TW, Lippod JC. Hydrogen induced crackingalong the fusion boundary of dissimilar metal welds. Weld J1999:31s–7s.

[27] Parker JD, Stratford GC. Review of factors affecting conditionassessment of nickel based transition joints. Sci TechnolWeld Joining 1999;4(1):29–39.

[28] Siewert TA. How to predict impact energy from stainless steelweld composition. Weld Des Fabr 1978;1:87–8.

[29] Ramazan K. Effect of solidification mode and morphology ofmicrostructure on the hydrogen content of duplex stainlesssteel weld metal. Mater Des 2004;25:1–9.

[30] Murti VSR, Sastry YR, Reddy GM. Some studies on the serviceweldability of HSLA steels using 18Cr, 8Ni, 6Mn typeaustenitic electrodes. J Mater Process Technol 1993;37:759–65.

[31] Pargeter RJ. Effects of arc energy, plate thickness and preheaton C–Mn steel metal hydrogen cracking. TWI Report no: 461,1992.

[32] Dayal RK, Parvathavarthini N. Hydrogen embitterment inpower plants. Sadhana 2003;28:431–51.

[33] Asami K, Sakai T. Hydrogen induced cracking at interfacebetween stainless steel overlay weld metal and base metal inpressure vessel. Trans Iron Steel Inst Jpn 1995;21(6):2–13.

[34] Matsuda F, Nakagawa H. Simulation test of disbondingbetween 2.25%Cr–1%Mo steel and overlaid austenitic stain-less steel by electrolytic hydrogen charging technique. TransJWRI 1984;13(1):159–61.

[35] Madhusudhan Reddy G, Mohandas T, Sarma DS. Coldcracking studies on low alloy steel weldments: effect offiller metal composition. Sci Technol Weld Joining2003;8(6):407–14.

[36] Bailey N. Welding of dissimilar metals. UK: The WeldingInstitute; 1996.

[37] Losz JMB, Challenger KD. HAZ microstructures in HSLA steelweldments. In: First United states—Japan symposium onadvances in welding metallurgy, Yokohama, Japan, 1999.p. 207–25.

[38] Kou S. Welding metallurgy. New York: Wiley; 1987.[39] Brand PC, De Keijser TH, Den Ouden G. Weld J 1993;72(3):93-

s–100-s.[40] Yurioka N, Suzuki H. Hydrogen assisted cracking in C–Mn

and low alloy steel weldments. Int Mater Rev1990;35(4):217–45.

[41] Hart PHM. Resistance to hydrogen cracking in steel weldmetals. Weld J 1986;65(1):14-s–22-s.

[42] Lee HW, Kang SW, Um DS. A study on transverse weld cracksin thick steel plate with the FCAW process. Weld J1998;77(12):503-s–10-s.

[43] Lee HW, Kang SW, Park JU. Fatigue strength, depending onposition of transverse cracks in FCAW process. Weld J2001;80(6):137-s–41-s.

[44] Madhusudhan Reddy G, Mohandas T. Effect of weldingparameters on residual stress distribution in flux cored arcwelding of high strength low alloy steels. In: NationalWelding Seminar, organized by Indian Institute of welding,India 1993. p. 33(a)–(d).

[45] Madhusudhan Reddy G, Mohandas T. Influence of weldingprocess and residual stresses on ballistic performance.J Mater Sci Lett 1996;15:1633–5.