-
A Study on Transverse Weld Cracks in Thick Steel Plate with the
FCAW Process
The transverse crack in thick plate welding is investigated
under simulated construction conditions
BY H. W. LEE, S. W. KANG AND D. S. UM
ABSTRACT. The transverse crack in thick plate welding is
discussed with respect to deposited metal. In recent years, many of
the new steel developments such as thermo-mechanical controlled
process (TMCP) have been intended to improve weldability. When TMCP
steel is used to achieve high strength with lean composition, the
weld metal is more likely to suffer hydrogen cracking than the
heat-affected zone (HAZ) of the base steel. Weld metal hydrogen
cracking is even more likely if alloying is necessary to match the
strength and toughness of the base metal. This is primarily due to
the more highly alloyed weld metal's increased susceptibility to
hydrogen cracking (Ref. 1).
One type of cold crack, referred to as a transverse crack, is
caused by the com- plex interaction of the diffusible hydrogen
supply, tensile residual stress and suscep- tible microstructure.
This form of cracking generally is not encountered when weld- ing
plate sections less than 10 mm thick. However, when thicker
sections (50 mm or more) are welded, welds are subjected to more
rapid cooling accompanied by more severe cooling stresses (Ref.
2).
Introduction
The various cracks that can occur in weld joints according to
welding condi- tions and processes are classified as "cold crack"
and "hot crack" according to occurrence temperatures.
Hot cracking, such as solidification cracks and liquation
cracks, are the most severe problems associated with
H. W. LEE is with the Welding Research Team of Samsung Heavy
Industries, Koje City, Korea. S. W. KANG and D. S. UM are with the
Research Institute of Mechanical Tech- nology, Pusan National
University, Korea.
the partially melted zone. The cause of hot cracking in the
partially melted zone is the combination of grain bound- ary
liquation and stresses induced by both solidification shrinkage and
ther- mal contraction during welding (Refs. 3, 4).
The transverse crack, a type of cold crack, occurs perpendicular
to the axis of the weld interface. It generally occurs at
temperatures below 200C (392F), ei- ther immediately upon cooling
or after a period of several hours. The time delay depends upon the
type of steel, the mag- nitude of the welding stresses and the hy-
drogen content of the weld (Refs. 5-7). However, most of the
literature on trans- verse cracks published thus far differs when
compared to the appearance of transverse cracks in actual
construction.
In this study, two EH 32 steel panels were welded to resemble
actual con- struction conditions. The appearance of transverse
cracks, hardness, impact, mi- crostructure and residual stresses
were then determined for two different weld- ing conditions.
KEY WORDS
Diffusible Hydrogen Intergranular (IG) Magnetic Particle
Inspection Microvoid Coalescence
(MVC) Quasi Cleavage (QC) Residual Stresses Stress Intensity
Factor Transverse Crack
Experimental Procedures
Test Panel
The size of the test panel was 2000 mm long x 1800 mm wide x 50
mm thick. The panel was fabricated from EH32 TMCP higher-strength
hull steel (as shown in Table 2), to provide test conditions
similar to actual construction conditions - - Fig. 1. To magnify
fabrica- tion-related weld residual stresses, the welding jig and
test panel were fillet- welded together.
Test Weldments
The specimen sections were welded in layers as shown in Fig. 2.
To compare the residual stresses and the position of occurrence of
the transverse cracks, the sections were welded under the follow-
ing conditions:
1) Below 30C (86F) of preheating and interpass temperatures.
2) Preheating and interpass tempera- tures of 100-120C
(212-248F).
The preheating temperature of 100C was obtained from the Yurioka
(Ref. 8) report shown in Fig. 3 (using Table 2, 50-mm-thick steel
plate, Ceq 0.34). The test specimens were welded at 100-120C in
consideration of ambient temperatures.
The panel was welded according to AWS A5.29 E8OT1-K2
specifications, using the flux cored arc welding (FCAW) process
(1.2 ~ diameter, electrode exten- sion of 25-30 mm); welding
parameters are shown in Table 1.
Chemical Composition/Strength
A spectroanalyzer was used to deter- mine the chemical
composition of the base and weld metal. Mean values of the three
specimens were then recorded in Table 2.
WELDING RESEARCH SUPPLEMENT [ 503-s
-
Fig. I - - Schematic diagram of weld panel.
I-
un i t : mm C A B 4 .
300 ~- 300 ~ 1,400
D
WELDING DIRECTION
4
A' ~B' C'
~,,~,,~,.---~ 3o" ~ /
----~l root opening: 10
(A - A')
I I
(B - B')
4#___ 3_~_
I I
(c - c )
Section A-A" : 1/3-specimen thickness weld deposit metal Section
B-B" : 2/3-specimen thickness weld deposit metal Section C-C" :
full thickness weld deposit metal
Fig. 2 - - Schematic diagram of weld deposit metal.
504-S J DECEMBER 1998
Hardness Traverses
Hardness was measured using the macro Vickers hardness test,
with a load of 5 kg and 10 s of loading time. Mea- surements were
made on transverse sections, 10 mm from the top surface.
Impact Test
The impact test was performed at 0, -20, -40 and -60C (32, -4,
-40, and -76F) using Charpy V-notch for deposited metal. The test
specimen loca- tion in the weldment is shown in Fig. 4.
Measurement of Residual Stresses
The surface residual stresses along the weld metal centerline (~
direction) were measured using the Rosette gauge hole- drilling
method after the specimen was cooled completely.
Diffusible Hydrogen Test
The diffusible hydrogen was mea- sured by glycerin method per
JIS Z3118. Before the hydrogen test, the steel plate was kept in
the furnace at 500C (932F) for 1 h and air-cooled to remove dif-
fusible hydrogen.
Distinction of Crack Position
To check the position and length of transverse cracks according
to changing preheating, interpass temperatures and welding layer,
the specimens were in- spected by ultrasonic testing. The surface
of the weld bead was then cut at 0.5-mm- depth intervals using a
milling machine and checked for accurate position and length of
transverse cracks using mag- netic particle inspection after each
ma- chining step.
Results and Discussion
Macro/Microstructure
The macrostructures of the weldments are shown in Fig. 5A and
B.
Some significant differences can be noted between the HAZ that
formed due to the welding pass and the HAZ located near the weld
interface. The grain size was very coarse in the HAZ near the weld
interface, resulting in the most brittle sec- tion of the weld
joints. Impact values im- proved in the reheated zone because the
grain boundary ferrite and Widmanstat- ten side plates were
transformed into pearlite and ferrite, and the grain size was
refined (Ref. 3).
Figures 6 and 7 are weldment mi- crostructures. Figure 6 is weld
joint "A" (preheating and interpass temperature
-
below 30C) and Fig. 7 is weld joint "B" (preheating and
interpass temperature 100-120C). Figures 6A and 7A show the
microstructure of the base metal, consisting of ferrite (white
area), pearlite (dark area) and bainite (slightly gray area). The
fine grain size results in ex- cellent strength and toughness (Ref.
9). The refined- and coarsened-grain re-
gions are shown in Figs. 6B and C and 7B and C,
respectively.
The refined-grain region was sub- jected to a peak temperature
just above the effective upper critical temperatures, Ac3, thus al
lowing austenite grains to nucleate. Such austenite grains decom-
posed into small pearlite and ferrite grains during subsequent
cooling. As
seen in Figs. 6B and 7B, the distribution of pearlite and
ferrite is not exactly uni- form because insufficient time was al-
lowed for the diffusion of carbon atoms due to the rapid heating
rate during weld- ing. The coarsened-grain region was sub- jected
to a peak temperature well above the Ac3 temperatures, thus
promoting the coarsening of austenite grains.
Table 1--Welding Parameters
Identification Welding Condition
A Preheating/interpass temperature below 30C
B preheating/interpass temperature 100-120C
Current Voltage Speed Heat Input Pass (A) (V) (cm/min)
(kJ/cm)
1 240-250 30 16 28 2-27 340-350 35 3741 26
1 240-250 30 15 29 2-27 340-350 35 38-42 25
Table 2--Chemical Composition of Base/Weld Metal
(%) C Si Mn P S Ni
0.18 0.10- 0.90- 0.040 0.040 0.40 EH32 TMCP max. 0.50 1.60 max.
max. max.
Base metal 0.09 0.38 1.35 0.015 0.005 0.03
Weld A 0.04 0.29 1.05 0.012 0.017 1.32 metal B 0.04 0.29 1.03
0.013 0.016 1.31
Mo V Ti TS YS El (kgf/mm 2) (kgf/mm 2) (%)
0.08 0.10 0.02 45-60 32.0 20.0 max. max. max.
0.02 0.002 0.02 52.8 38.0 31.0
0.02 0.017 0.01 69.4 63.7 22.8 0.02 0.018 0.01 66.3 61.4
23.4
250 (~c)
~ 150
~ 100
H.~ = 5=~/100~ WM -- H.I. = 1.7Klknm
Ambient ~ = 10"(2
, / 0 ~'75 75 60 50 40 30 25 20 15 S. lOrnm
0.2 0.3 0.4 0.5
Carbon equivalent, Ceq
0.6
1 1 1! ,2 ,3 4 ,5,6 7 8
Fig. 3 - - Diagram of preheating temperature for Ceq and steel
plate thickness.
5O
10mm
Fig. 4 - - The position of Charpy V-notch impact test
specimen.
Fig. 5 - - Macrostructure of weld joint near section C - C'. A -
- Preheating/interpass temperature below 30C; B - - preheat-
ing/interpass temperature 100-120C.
WELDING RESEARCH SUPPLEMENT I 505-$
-
w
Hi @'j isu
m q
~L ,G I W
W
Wil l
0
Fig. 6 - - Microstructure of weld jo int A (preheating and
interpass temperature below 300 . A - - Base metal; B - -
grain-refined zone; C - - grain-coarsened zone; D - - weld
metal.
Fig. 7 - Microstructure of weld jo int B (preheating and
interpass temperature 100-1200 . A - - Base metal; B - -
grain-refined zone; C - - grain-coarsened zone; D - - weld
metal.
Because of the relatively high cooling rate and the large grain
size in this region, acicular ferrite rather than blocky ferrite
formed at grain boundaries-- Figs. 6C and 7C. Figure 7D is an
optical micrograph taken from the deposited weld metal area
revealing grain boundary ferrite, Wid- manstatten ferrite and
acicular ferrite. To improve mechanical properties such as tensile
and toughness, acicular ferrite has
to form fully instead of grain boundary fer- rite and
Widmanstatten ferrite.
More amounts of acicular ferrite can be observed in Fig. 7D when
compared to Fig. 6D.
Residual Stresses of Weld Joints
Welding induces high residual stresses in the vicinity of the
weld. The
residual stress is caused by restraining the free contraction of
the thermoplastically deformed weld zone during weld cool- ing.
Therefore, the welding residual stresses are sometimes referred to
as re- straint stresses. A geometrical notch in the weld joint
further induces local stress concentration. In most cases,
hydrogen- assisted cracking is initiated at a notch of the weld
made under restraint.
The residual stresses measured at the surface of a deposited
metal in a longitu- dinal direction of weld interface are shown in
Fig. 8. In all measured points, the residual stress values for a
specimen- welded preheating and interpass temper- ature below 30C
was higher than the preheating and interpass temperature of
100-120C. Transverse crack occurrences are caused by the hardness
of deposited metal, diffusible hydrogen contents and tensile
residual stresses in the longitudinal direction of the weld
interface.
Diffusible Hydrogen Contents
Hydrogen-assisted cracking is a severe problem in the welding of
thick steel plate that occurs when the follow- ing three factors
are simultaneously pre- sent: diffusible hydrogen in weld metal,
high stress and susceptible microstruc- tures. The hydrogen
dissolved in a weld metal is proportional to the square root of the
partial pressure of the hydrogen gas.
The following sources of weld metal hydrogen are considered in
FCAW (Ref. 10):
1) Moisture in flux 2) Moisture in CO 2 gas 3) Organic substance
in flux 4) Hydrogen in wire steel and steel
plate 5) Moisture in atmosphere 6) Extraneous hydrogenous
material,
e.g., moisture, grease and paint Hydrogen dissolved in a steel
matrix
is diffusible, thereby causing hydrogen embrittlement. The weld
metal hydrogen content is generally expressed by the content of
diffusible hydrogen. The three methods of measuring diffusible
hydro- gen contents are:
1) Glycerin method (Hjl s) 2) Mercury method (Hi] w) 3) Gas
chromatograph method (HG_ c) The test results of these three
methods
are related as follows (Ref. 11):
HHw = 1.27Hji s + 2.19
HG_ C = 2HjI s + 0.3
where Hil w, HG_ C and HjI S are the weld metal diffusible
hydrogen content per 100 g of deposited weld metal.
The hydrogen contents, which de-
506-s I DECEMBER 1998
-
(A )
7O
50
~ 10 i
ol
---O-- preheating/interpass temp. below 30"C ] - " "0"" preheat
ing / in terpass temp. 100-120"C J
I I I I [ I I A B C D E F G
(B)
Fig. 8 - - Distr ibutions o f surface residual stress for (~
direction in deposited metal. A - - Position of attached Rosette
gauge; B - - results of surface residual stresses.
800
s fl,O0
m 0 O. at "U 400
g Q
~ 200
'I"
o,oo
240A X 30V, 25CPM
(~ 350A X 35V, 25CPM
1 0 ~) (D
/ [ I I I I 2 4 18 48 72
Exposure Time (Hour)
Fig. 9 - - Hydrogen content profiles depending on welding
conditions.
v
-
UJ
0 u~
400 --
300
200
100
preheat/interpass temperature 100~120"C
preheat/interpass temperature below 300C
-80
!
t
I I I I I -60 -40 -20 0 20
Temperature ( C)
Fig. I 0 - - Results of Charpy V-notch impact tests for weld
metal.
24000- ~ prebeating/interpass temp. 100-120"(."
preheating/interpass temp. below 30"C
220.00
(~ 200.00
180.00
'10
~ 16000
W
14000
-soo -4oo ooo 4.oo 8oo ~2oo Distance from the weld interface
(ram)
Fig. 11 Hardness traverses 10 mm from top surface.
pend on welding conditions, were mea- sured by the glycerin
method and are shown in Fig. 9. These data indicate that welds made
with the FCAW electrode have hydrogen contents of approx- imately
3-4 mL. Most diffusible hydrogen escaped within 2 h after weld- ing
as shown in Fig. 9. However, when welding conditions were changed,
there were no significant hydrogen contents.
Impact Properties
Figure 10 shows the Charpy V-notch impact test results for weld
metal. Ab-
sorbed energy of preheating and inter- pass temperatures of
100-120C in weld joint B are higher than preheating and interpass
temperatures of 30C in weld joint A, due to higher cooling
rate.
Hardness Traverses
Figure 11 shows the hardness tra- verses 10 mm away from the
weld sur- face. When the preheating/interpass temperature is below
30C, the value of hardness (HV) in deposited weld metal is 10-15
higher due to the rapid cooling rate. The hardness of the weld
metal
depends on the preheating/interpass temperature, and when the
preheating/ interpass temperature is low, the weld metal becomes
more susceptible to transverse crack.
Distinction of Crack Position
In weld joint B, when preheating and interpass temperatures were
100-120C, no transverse cracks were detected. However, transverse
cracks were de- tected for the specimen welded with pre- heating
and interpass temperatures below 30C in weld joint A.
WELDING RESEARCH SUPPLEMENT I 507-s
-
(A) I(C) 16.0 mm
Fig. 12 - - MPI results of transverse cracks for 35-mm weld
joint A (depth below weld top surface is shown at top right-hand
corner).
(A) 9.0 mm
| lJi~ I I I lliIIIII [I
}l[,l'il[!ll}flllLllIi[llillIItillIl!iiill ~ ' ~a ~ 9o INn
, ; 2
(B) 9.5mm I(C)
I
11.0 mm
(r- B I I
~ ^ nm
(G (H) 28.0 mm
Fig. 13 - - MPI results of transverse cracks for 50-mm weld
joint A (depth below weld top surface is shown at top right-hand
corner).
goR-~ I IOFCFMRER 1998
-
It was also noted that the number of transverse cracks increased
as the weld- ing layers were increased. No cracks were observed in
the one-third complete sample. However, transverse cracks were
formed in specimens welded in two- thirds of their thickness (35
mm) and the full thickness of the weld joint (50 mm).
Figure 12 shows the morphology of detected transverse cracks for
specimens welded in two-thirds of their thickness (35 mm) at
various depths from the weld surface. Figure 12A shows the morphol-
ogy of the transverse cracks that ap- peared for the first time,
located 9.5 mm in depth from the weld surface. Some of these cracks
can also be seen in Fig. 12B, which also shows the formation of new
cracks at a depth of 10.0 mm from the weld surface. These cracks
completely disappear and two new sets of cracks are visible in
Fig.12C at a depth of 16.0 mm from the weld surface. The cracks
disap- pear as the distance from the weld sur- face is increased
and a new set of cracks is formed - - Fig. 12D.
Figure 12E shows the continuation of cracks detected in Fig.
12D; however, their number has decreased consider- ably. Transverse
cracks completely dis- appear at a depth of 20.0 mm from the weld
surface-- Fig. 12F.
It can be seen from the macrographs of Fig. 12 that transverse
cracks have a pattern of appearing and disappearing in locations at
depths 9.5-17.0 mm away from the weld surface.
The morphology of transverse cracks in the full-thickness joint
is shown in Fig. 13. Cracks at a depth of 9.0 and 9.5 mm from the
weld surface are shown in Fig. 13A and B. A comparison of these two
macrographs shows that transverse cracks have a clearer morphology
in Fig. 13B. Additionally, the morphology of cracks in Fig. 13E and
F is more delicate, with a number of small-size cracks located in
the surrounding area of the larger cracks than the cracks detected
in Fig. 13C and D.
In the full-thickness weld, the trans- verse cracks also have a
pattern of appearing and disappearing at a depth between 9.0-27.0
mm away from the weld surface and completely disappear at a depth
of 28.0 mm - - Fig. 13H.
The position of transverse cracks rela- tive to the weld layer
is shown in Fig. 14. For the specimen welded at two-thirds
thickness (35 mm), transverse cracks were located between weld
layers 4-6, as shown in Fig.14A. For the specimens welded at
full-thickness (50 mm), trans- verse cracks were located between
weld layers 5-8, as shown in Fig. 14B.
Takahashi, etal . (Ref. 12), have shown that transverse cracks
are initiated in the
weld metal just / below the final layer of welds and 50 are
gradually prop- agated toward both the top and bottom surface. This
oc- curs because the largest residual stresses and the highest
concentra- / tion of diffusible / hydrogen contents t5,.0 can be
found in these locations.
In this study, the position of trans- verse cracks differed from
that of the Takahashi report, since cracks were detected at a con-
stant distance from the top of the weld surface (e.g., 9.5-10.0 mm,
and not just below the final layer as reported by Takahashi). This
is due to a large restraint stress under actual construction condi-
tions, as compared to a small test piece.
Figures 15A and B show an optical micrograph of the middle and
edge of transverse cracks. The formation of these cracks did not
follow the grain boundary ferrite; rather, they propagated across
the grains.
From fracture morphology, it is noted that transverse cracks
occur in high stresses. Microscopic fracture modes from the Beachem
(Ref. 13) report are shown in Fig. 16. These illustrations show the
tip of cracks growing from left to right under four different K
(stress intensity fac- tor) conditions, with the K decreasing from
Fig. 16A through D. This represents a suggested explanation of the
changes in the observed fracture modes.
Conclusions Macrostructure appearance of the
6 layer~~ t 4 layer'---~ ~ --3.- 9.5-20.0 I I
(A )
I ~ I unit :mm
(B)
Fig. 14 - - The transverse crack position according to changing
welding layer. A - - 35-rnm weld joint; B - - 50-ram weld
joint.
transverse crack and mechanical proper- ties such as hardness,
impact and resid- ual stress measurement was studied for EH 32 TMCP
50-mm-thick plate welded with FCAW under the condition of changing
preheat and interpass tempera- ture. The results of this study can
be sum- marized as follows:
1) Transverse cracks were detected in the specimen welded with
preheat- ing and interpass temperatures below 30C, but cracks were
not detected for the specimen welded with preheat and interpass
temperatures of 100-120C.
2) Two different locations of crack formation were detected in
this experi- ment as follows:
a) In the specimen welded at two- thirds thickness of the joint,
cracks were initiated at a distance of 9.5-10 mm away from the top
of the welded surface, be- tween layers 4-6.
b) In the specimen welded at full thickness, cracks were
initiated at a dis-
5 0 ~ : ;4~2i~; ' .~ " aO~m i
Fig. 15 - - Optical microstructure o f transverse cracks. A - -
Middle of crack; B - - crack edge.
WELDING RESEARCH SUPPLEMENT I 509-s
-
(A)
(c)
(e)
/ . . . . . . . . . . . \
(D)
Fig. 16 - - Microscropic fracture modes. A - - MVC with high
stress intensity factor; B - - QC with intermediate stress
intensity factor; C - - IG cracking with low stress intensity
factor; D - - IG cracking with assisted hydrogen pressure.
tance of 9.5-10.0 mm away from the top of the welded surface,
between layers 5-8.
3) Hardness values of preheating and interpass temperatures
below 30C were higher than preheating and interpass temperatures
below 100-120C in de- posited weld metal.
4) The residual stress values for the specimen welded with a
preheating and
interpass temperature below 30C was higher than the specimen
welded with a preheating and interpass temperature at 100-120C.
5) The result of weld metal impact test shows higher impact
values for the weld condition with preheating and interpass
temperature at 100-120C, compared to the specimen welded with
510-s I DECEMBER 1998
a preheating and interpass temperature below 30C.
References
1. Bailey, N., and Wright, M. D. 1993. Weldability of high
strength steels. Welding and Metal Fabrication, pp. 389-396.
2. Metals Handbook. 1973.9th ed., Vol. 6 ASM International,
Materials Park, Ohio, pp. 129-130.
3. Kou, S. 1987. Welding Metallurgy, John Wiley and Sons, New
York, N.Y., pp. 249, 326.
4. Welding Handbook, 1987. 8th ed., Vol. 1. American Welding
Society, Miami, Fla., pp. 230-231.
5. Signes, E. G., and Howe, P. 1988. Hydrogen-assisted cracking
in high-strength pipeline steel. Welding Journal 67(8): 163-s to
170-s.
6. Suzuki, H. 1978. Cold cracking and its prevention in steel
welding. Transactions of the Japan Welding Society, pp. 82-86.
7. Hart, P. H. M. 1986. Resistance to hydrogen cracking in steel
weld metals, weld- ing Journal 65(1 ): 14-s to 22-s.
8. Yurioka, N. 1995. A chart method to determine necessary
preheat temperature in steel welding. Journal of Japan Welding
Society, pp. 347-350.
9. Lee, H. W., and Kang, S. W. 1996. A study on microstructure
and thoroughness of electrogas weldments. Journal of the Korea
Welding Society, pp. 68-74.
10. Yurioka, N., and Suzuki, H. 1990. Hydrogen assisted cracking
in C-Mn and low alloy steel weldments. International Materials
Review, pp. 217-249.
11. JlS Z3118 Method of Measurement for Hydrogen Evolved from
Steel. 1986. Japanese Standard Association.
12. Takahashi, E. 1979. Relations be- tween occurrence of the
transverse. Journal of Japan Welding Society, pp. 855-872.
13. Beachem, C. D. 1972. A new model for hydrogen-assisted
cracking. Metallurgical Transactions 3(2): 437-451.