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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
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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 [email protected]
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.
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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
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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
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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
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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.
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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
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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
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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.
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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).
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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.
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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
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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.
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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
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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 havesignificant 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) exhibitedgreater 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.
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