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Effects of Shielding with Various Hydrogen-Argon Mixtures on
Supermartensitic Stainless Steel TIG Welds
Article in Materials Testing · May 2010
DOI: 10.3139/120.110135
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Effects of Shielding with Various
Hydrogen-Argon Mixtures on
Supermartensitic Stainless Steel
TIG Welds
306
Dan Eliezer, Y. Nissim,
Beer-Sheva, Israel, and
Thomas Kannengießer, Berlin
A number of common defects in stainless steel welding result from
the presence of hydrogen in the weld. The service life of the stainless steel joints is further significantly dependent on the presence
of hydrogen in the respective environment and the susceptibility
of various weld microstructures to hydrogen degradation.
As a relatively new materials generation, supermartensitic stainless steels (SMSS) are increasingly applied to substitute more
expensive alloys, particularly in the oil and gas industries. As a
result of their martensitic microstructure these alloys are prone to
hydrogen assisted cracking (HAC). The resistance of SMSS to
hydrogen assisted stress corrosion cracking (HASCC) during sour
service has been extensively studied, predominantly for industrial
purposes. Studies are primarily conducted with parent materials
based on standard test procedures. The principal hydrogen behavior in welded SMSS microstructures has been less investigated.
The central objectives of the study are to determine the hydrogen
interactions with the microstructure of a Gas Tungsten Arc (GTA)
welded SMSS and hydrogen trapping mechanisms.
The interactions of hydrogen with various Tungsten Inert Gas (TIG)
welded SMSS microstructures are investigated by means of X-ray
diffraction (XRD) and optic (OM) and electronic microscopy (SEM).
A number of methods have been employed for the estimation of the
quantities of absorbed hydrogen. Hydrogen interaction with structural defects and the characteristics of hydrogen desorption have
been studied by means of thermal desorption spectroscopy (TDS),
and hydrogen content measurements (LECO analyses). The effects
of the respective microstructure on hydrogen absorption and
desorption behavior are discussed in detail.
Supermartensitic stainless steels are
essentially iron-based alloys which contain chromium, nickel, and molybdenum. They contain approximately the
same amount of chromium as conventional martensitic stainless steels of
the AISI 410 and 420 series, but their
carbon content is reduced. The addition
of up to 6.5 wt.-% nickel substantially
decreases the martensite transformation towards lower temperatures; weld
microstructures may contain considerable amounts of retained and annealing
austenite.
Three different levels of nickel and
molybdenum, which were commercially
available, were therefore selected for fitness purposes. Lean alloys are intended
to be used in sweet service and contain
nickel below 2 wt.-% and molybdenum
below 1 wt.-%, and are fabricated as low
© Carl Hanser Verlag, München MP Materials Testing 52 (2010) 5
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carbon (LC) steels with carbon contents
below 0.02 wt.-% and with nickel and
molybdenum amounts of between 2 and
4 wt.-% and of approximately 1 wt.-%,
respectively, medium alloys have been
developed for intermediate sour service
applications. At increasing H2S levels
and decreasing pH extra low carbon
(ELC) fat grades are applied. These
materials contain even less carbon, i. e.
under 0.01 wt.-%, but even more nickel,
up to 6.5 wt.-%, and more molybdenum
up to 2.5 wt.-% [1-4]. The base materials
are usually heat treated in several stages
and consist of considerable portions of
annealing and retained austenite, depending on the nickel content [2].
Supermartensitic alloys have been
commercially available since the early
1990’s and have been increasingly applied, almost exclusively in the oil & gas
industries. In these industries the selection of materials which fit exactly
with the intended applications has become increasingly important for economic reasons. For example, supermartensitic stainless steels represent an attractive cost reducing alternative to higher
alloys and more expensive materials,
and are thus increasingly applied as
flowline materials at several North Sea
oil and gas fields. At these locations,
welded components are subjected to
sour service conditions, providing a potential risk for hydrogen uptake, degradation and cracking of supermartensitic
stainless steels.
As another potential failure risk for
welded SMSS, Hydrogen Assisted Cold
Cracking (HACC) has to be avoided during welding. Even small hydrogen concentrations, trapped inside the material,
might lead to failure. Consequently, a
comprehensive knowledge of hydrogentrapping interactions is necessary to
make any decision and/or judgment as
to whether a trap site or a particular
trapped hydrogen content is detrimental
to the safe service operation of welded
joints of such materials.
Several industrial application-oriented studies have been performed to investigate the sour service and respective
HASCC resistance of SMSS, predominantly in the as-delivered, but also in
the welded condition. Limited investigations have further been undertaken to
evaluate their cold cracking resistance.
A preliminary comparison revealed the
same fracture topography SMSS weld
metals for HACC and for HASCC [1].
52 (2010) 5
C
S
P
Mn
Si
Ni
Cr
Mo
N
0.06
0.0009
0.021
1.870
0.294
6.498
11.65
2.330
0.009
MP
Table 1. Chemical composition of the SMSS (wt.-%)
Therefore, the objective of the present
study is to investigate hydrogen interactions with GTA welded supermartensitic stainless steel microstructures in
greater detail. Particularly with respect
to adsorption and desorption characteristics, as well as to hydrogen trapping
mechanisms.
Experimental Approach
The chemical composition of the tested
supermartensitic stainless steel is listed
in Table 1. In order to investigate hydrogen affects on the microstructure, as
well as on the absorption and desorption
behavior of hydrogen in supermartensitic stainless steel, the alloy was exposed to hydrogen in three different
modes. These are as follows: electrochemical hydrogenation, gaseous-phase
hydrogenation and hydrogen introduced
by GTA welding via a mixed Ar + H2
shielding gas.
Hydrogen charging was carried out
electrochemically (cathodic charging) at
room temperature. The cathodic charging technique generates H+ on the sample’s surface. A power supply (galvanostat) impresses a constant current on
the electrolytic cell, such that the specimen is cathodic relative to an inert electrode such as platinum (Pt). The anode,
a platinum wire, is located symmetrically to the cathodic polarized specimen,
thereby uniformly distributing the potential around the specimen. The specimens were cathodically hydrogenated at
identical charging times, using a current
density of 50 mA/cm2. The charging was
carried out in 1N D2SO4 + 0.25 g NaAsO2
electrolyte.
Gaseous phase charging was carried
out in H2 UHP 99.99 atmospheres, using
a pressure of 1 atm during 6.5 hours at
room temperature and 650 °C.
The charging of the weld metal with
hydrogen was performed by a Gas Tungsten Arc Welding (GTAW) using the
parameters summarized in Table 2.
Prior to the electrochemical and gas
phase hydrogen exposure, samples were
cut from SMSS plates, and were mechanically polished up to 0.05 μ. The
total content of hydrogen/deuterium
absorbed in the alloy was measured by a
LECO RH-404 hydrogen determination
system, using the test protocol described in the previous literature [5]. Hydrogen/Deuterium evolution and trapping
characteristics were studied by thermal
desorption spectroscopy (TDS), a technique that involves measurement of the
desorption rate of gas atoms, soluted or
trapped in the material, while heating
the sample at a known rate [6-7]. After
ultrasonic cleaning in ethanol, the deuterium charged specimens were placed
in the specimen holder; the system was
sealed and pumped down to 10 mPa.
The heating rates (ramps) were 3, 5, and
7 K min-1, and the temperature range
was between 25-600 °C, both parameters being programmed into a temperature controller. While heating, the mass
spectrometer was placed in a continuous
mode for scanning atomic masses in the
range of between 3.5-4.5 amu.
Shielding Gas
Type
Shielding Gas
[l/min]
Current
[A]
Voltage
[V]
Weld Speed
[cm/min]
Argon + 0% hydrogen
10
145
10
30
Argon + 2% hydrogen
10
145
11.5
32
Argon + 5% hydrogen
10
145
11.5
32
Argon + 7.5% hydrogen
10
145
11.5
39
Table 2. Welding parameters of SMSS
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Figure 1. Hydrogen contents in different welding areas of SMSS TIG welded in gas shield atmospheres containing different hydrogen percentages
Specimens for optical microscopy
were prepared by mechanical polishing
through 0.05 μm alumina. The samples
were finally etched using one of the following etchings:
1. 5 gr. cupric chloride in 100 cc hydrochloric acid + 100 cc ethyl alcohol +
100 cc water
2. 10 gr. ferric chloride in 20 cc hydrochloric acid + 4 cc nitric acid.
The use of different etchings was necessary to highlight specific features and
phases and thus obtain complementary
information on the microstructures.
Digital photographs were taken using a
Leica DMR digital camera. The specimen’s surface was observed using a
JEOL JSM 5600 scanning electron micro-
scope (SEM). The machine was operated
at an accelerating voltage of 15 kV.
The microstructure and phase characterization of the above specimens were
studied by means of X-Ray diffraction
(XRD) analyses, using a Rigaku Type
2000 X-ray powder diffractometer with
a CuKα radiation (λ = 1.54059 Å) and
a graphite monochromator diffracted
beam. The data was collected in an angular range of 41° < 2θ < 54°, 61° < 2θ < 94°
by steps of 0.05° (2θ) with a constant
counting time of 4s by step.
Results and Discussion
Hydrogen Absorption. The amount of
hydrogen absorbed in the alloy in the
Sample
As-received
4.06 ± 0.28
Charged for 24 h with deuterium
165 ± 11.55
Gas phase hydrogenation (R.T., 6.5 hours, 1 atm)
2.16 ± 0.11
Gas phase hydrogenation (650 °C, 6.5 hours, 1 atm)
1.87 ± 0.15
TIG welded in Ar 0 % H2
6.7 ± 0.34
TIG welded in Ar 2 % H2
6.2 ± 0.42
TIG welded in Ar 5 % H2
5.83 ± 0.25
TIG welded in Ar 7.5 % H2
8.64 ± 0.60
Table 3. Hydrogen content in SMSS alloy
308
Absorbed Hydrogen
(ppm)
as-received state and after exposure to
hydrogen was measured by means of a
LECO analysis. Results are summarized
in Table 3. The hydrogen contents in different welding areas of the TIG welded
specimens welded with gas shield atmospheres containing different hydrogen
percentages are presented in Figure 1.
It needs to be noted that in its asdelivered condition the base material
contains about 4 ppm.
Table 3 demonstrates the different
ways by which the hydrogen was introduced to the alloy significantly influence
the amount of absorbed hydrogen.
It can be generally stated that the
higher the fugacity of hydrogen in the
atmosphere the larger the hydrogen concentrations absorbed in the alloy. The
hydrogen content absorbed in the electrochemically charged specimen is the
highest and differs at a rate of about two
orders of magnitude from that of the
as-received specimen. This result can be
attributed to the high hydrogen concentrations produced on the sample surface,
suppressing most of the recombination
during the electrochemical charging
period, as reported in former studies [2].
The distribution of hydrogen in the volume is dependent upon the hydrogen
charging technique, i. e. if hydrogen is
introduced from the surface, as occurs
during electrolytic or gaseous-phase
charging, or is introduced by a welding
arc and homogenously distributed, as
occurs in a weld metal. The internal hydrogen distribution is also significantly
dependent on the various weld microstructures, due to their different diffusivities and solubility rates [2, 8].
It should be noted that the hydrogen
amount absorbed during gas-phase
hydrogenation is considerably low, and
even lower than in the as-delivered base
material. This is to be attributed to prior
degassing in the process. From the
above mentioned result, which shows
some consistency with a similar experimental study with carbon mild steels
[9-10], it can tentatively be concluded
that at low pressures hydrogen will not
enter the material from a gaseous phase
even at elevated temperatures of up to
650 °C. This is as long as the material is
not considerably plastically deformed
under a static or dynamic mechanical
load.
From Figure 1, it can be seen that the
different weld microstructures significantly influence the amount of hydrogen
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absorbed in the specimen, due to their
variable solubility rates. The lowest concentrations have been measured in the
heat affected zones (HAZ) which are to
be attributed to the heat introduced into
the HAZ during welding. This result
stands in contrast to a recent investigation of the HAZ multilayer GTA welded
SMSS of similar chemical composition
where the highest or at least the similarly high hydrogen concentrations have
been found in the HAZ as compared
to the heat affected zone. As shown in
Figure 2, such high hydrogen levels in
the HAZ probably have to be attributed
to a high fraction of retained or annealing austenite in the HAZ of about 35 %.
Such annealing and austenite forming
effects are not present in the mono-layer
weld investigated here, and therefore
less hydrogen is accumulated in the HAZ.
Also, in contrast to the previous study
with multilayer welds [2], somewhat
higher hydrogen concentrations than in
the HAZ have been found in the weld
metal of the mono-layer weld investigated. The hydrogen content in the weld
metal exceeds that of the base material
only if very high hydrogen percentages
are added in the shielding gas, i. e. at a
fraction of 7.5 % in argon.
Hydrogen’s Effects on the Microstructure. The exposure parameters
extensively influence the amount, quality and depth of hydrogen absorption.
This means that as a consequence of the
hydrogenation process different microstructural morphologies might be obtained. Also, the respective phase transformations are dependent on the amount
and way in which hydrogen is introduced into the material. For instance,
hydrogen represents an interstitial, and
might shift the martensite formation towards lower temperatures during the
cooling period after welding.
The initial microstructure and XRD
pattern of the studied material are
shown in Figures 3 and 4. The phase
analysis of SMSS XRD patterns reveals
the existence of two phases of BCCmartensite and FCC-austenite arranged
in a lamellar structure.
The γ-phase reflections exhibits significant broadening, but the FCC lattice
parameter remains unchanged.
The α-phase reflections, on the other
hand, exhibits negligible broadening
and the BCC lattice parameter contracts
by approximately 0.8 %. This result can
be attributed to the fact that hydrogen
52 (2010) 5
MP
Figure 2. Solubility evaluated by the sub-surface concentration of hydrogen dependent on the
austenite content in the weld microstructures of similar nickel and molybdenum alloyed SMSS
multi-layer steel welds [2]
Figure 3. X-Ray diffraction patterns of electrochemically charged SMSS alloy with deuterium for
24 hours
Figure 4. Typical microstructure of the as-received SMSS, a) optical micrograph, b) SEM micrograph
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Figure 5. X-Ray diffraction patterns of TIG welds in gas shield atmospheres containing Ar +7.5 %
hydrogen, from different welding areas
Figure 6. Optical micrographs of the SMSS TIG weld microstructures shielded with Ar +7.5 %
hydrogen, a) weld metal and HAZ at the fusion line, b) HAZ, c) weld metal, d) weld metal
has very different diffusion rates in
austenitic and in martensitic microstructures. The rate is higher in body
centered cubic lattices, and remains
high even at low temperatures. Therefore, the retained and continuous γ
phase in the fully lamellar microstructure of the SMSS alloy offers a blocking
site for hydrogen diffusion.
The microstructure and XRD pattern
of the different weld microstructures are
shown in the Figures 5 and 6 above. The
310
phase analysis of SMSS XRD patterns
reveals the existence of two phases BCCmartensite and FCC-austenite according
to the WRC 1992 diagram [11]. As to be
expected from such a mono-layer weld,
the base material contains a larger
amount of retained or annealing austenite in comparison to the HAZ area, while
the weld metal completely transformed
into martensite.
Figure 7 exhibits X-Ray diffraction
patterns of TIG welded SMSS alloys,
from the weld metal, at different hydrogen percentages in the atmosphere. As
the hydrogen percentage in the atmosphere increases, more hydrogen is introduced into the completely transformed
martensitic lattice and consequently,
the α-phase reflections exhibits significant broadening.
Results of XRD patterns after gas-phase hydrogenation are shown in Figure 8.
Evidently the soluted hydrogen concentrations under these charging conditions
are too small to cause any broadening of
the γ- and α-phase reflections, i. e. no
lattice parameter expansion occurs at
such conditions.
Desorption Charcteristics. In contrast to conventional hydrogen measurement and extraction technologies used
for welding applications, such as the
mercury procedure or carrier gas hot
extraction, the thermal desorption spectroscopy (TDS) technique represents a
much more suitable, reliable, and expressive method to characterize the
hydrogen evolution process, and to assess hydrogen trapping characteristics
in a microstructure. For this reason it
has been utilized within this study.
The spectra analyses were supported by
data obtained from a variety of other
experimental techniques, such as LECO
hydrogen quantity analyses, XRD, and
microstructure investigations by means
of optic and electronic microscopy [5-6].
The trapping phenomena in steels are
not well understood and their connection with hydrogen assisted cracking
is unclear. For instance, it is unclear
whether hydrogen deeply trapped
within a Titanium Carbide may not be
(re-)activated during the straining of
the microstructure thereby contributing
to cracking. Furthermore, although it
has been shown and it can be anticipated that hydrogen solubility increases
and its diffusivity decreases with the
increase of chromium carbide precipitation in SMSS [2] it is not evident that
these hydrogen accumulations contribute to inter-granular cracking [1]. At
the same rate, it is however difficult to
obtain unambiguous experimental information on the influence of traps for
cracking, partly as a result of the small
hydrogen concentration involved. Yet
there is a generally consensus in past
literature that the lattice diffusivity of
hydrogen in steel is strongly retarded by
the presence of microstructural inhomogeneities including grain boundaries,
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dislocations, carbides, and nonmetallic
particles, all acting as potential hydrogen trapping sites [12-14].
The capacity of a microstructural feature to act as a trapping site is governed
by the potential-energy well (and binding energy (Eb)) relative to the normal
interstitial positions in the Fe-lattice.
With the increasing strength of the trapping site relative to the energy well
depth, i.e. the peak height of the free
energy graph, traps are categorized as
reversible and then irreversible. In the
absence of traps, hydrogen diffusion
occurs by a random walk between interstitial lattice positions, usually with a
very short residential time for the diffusing of a hydrogen atom. Hydrogen
diffusion can be described by the ideal
lattice diffusion coefficient, DL, and
governed by the activation energy for
diffusion associated with the energy
barrier (on lattice migration energy
(Em)) between such interstitial positions
E
––m––
[13], i. e. DL = D0 · e––R·T
.
The activation energy in this study is
calculated by the Lee and Lee model [7],
which is based on the following assumptions: (a) the controlling process is a
first order, which is described by detrapping and diffusion to the surface, i. e. the
chemical absorption energy (Echem) is
smaller than the migration energy
(Em) or smaller than the binding energy
(Eb), and second order desorption phenomena, such as molecular desorption
accompanied by recombination of adsorbed atoms, are disregarded (b) with
the increase of the heating rate of the
specimens the peaks are shifted to
higher temperatures.
Taking into account the shift of the
peaks at different heating rates, the
effective energy of hydrogen desorption
was evaluated from the slope.
Characteristics of hydrogen desorption from the SMSS specimens are presented in Figures 9 to 11. The calculated
results from these TDS graphs are summarized in Tables 4 to 6. For all of the
SMSS specimens subjected to hydrogen
in the “as- received”, TIG welded, and
gas phase hydrogenated condition, the
TDS spectra exhibit a common feature
that initial degassing occurs until the
vacuum reaches 10-7 to 10-8 mbar. Furthermore, all the spectra exhibited a
first peak at relatively low temperatures,
i. e. around room temperature, indicating a larger loss of hydrogen. These
peaks are probably associated with hy-
52 (2010) 5
MP
Figure 7. X-Ray diffraction patterns in the base material of the TIG welded SMSS at different
hydrogen percentages in the atmosphere
drogen escaping from the specimen’s
surface at a relatively low activation
energy of approximately 4.5 kJ.
A second evolution peak occurs in all
spectra over a wide temperature range
from about 80 °C up to 200 °C with an
activation energy of approximately
5.7-6.7 kJ.
The spectra for the as-received material (Figure 9a), for the 1 atm gaseouscharged material (Figure 9b), for the
weld metal, and the base material produced without hydrogen in the shielding
gas (Figure 10a) and finally for the
base material of the weld shielded with
7,5 vol.-% hydrogen in the argon atmosphere (Figure 10b) – all have nearly
identical shapes. As to be expected,
no additional hydrogen was taken up in
the material during welding without
hydrogen. Also, no hydrogen was traced
in the base material during welding with
7.5 vol.-% in the shielding gas. Most
importantly, almost no additional hydrogen was taken up during charging
at room-temperature in a gaseous hydrogen environment.
However, concerning the 7.5 vol.-%
hydrogen shielded weld metal and after
gas-phase hydrogenation at high temperature, it was observed that hydrogen
evolution reaches a characteristic peak
at approximately 285-300 °C, and covers a wide temperature range between
Figure 8. X-Ray diffraction patterns (1 atm, 6.5 hours) of SMSS after gas-phase hydrogenation,
charged at room temperature and 650 °C
311
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a)
b)
Figure 9. TDS spectrum of SMSS at different charging temperatures and at a heating rate of 3 °C/min a) as-received material, b) gas-phase charged
(1 atm, 6.5 hours)
a)
b)
Figure 10. TDS spectrum for different SMSS welded microstructures of alloys, TIG welded in different shielding gas atmospheres at a heating rate of
3 °C/min, a) Ar +0 % H2 , b) Ar +7.5 % H2
a)
b)
Figure 11. TDS spectrum of samples electrochemically charged with deuterium for 24 hours, TIG welded in gas shield atmospheres containing
Ar +7.5 % H2 at a heating rate of 3 °C/min, a) as-received material, b) weld metal
ΔT = 140 °C and 360 °C with an activation energy of approximately 8.5 to
9.1 kJ. In comparison to the literature,
the activation energy values are calculated close to migration energy: pure annealed Fe, Em is 7 kJ/mol (0.07 eV/atom),
and DL is 1.3 · 10-5 cm2/s at 25 °C [2].
312
The specimens subjected to electrochemical hydrogen charging exhibited
completely different hydrogen evolution
characteristics. In particular, a large
peak was observed at approximately
150 °C that covers a wide range of temperature between ΔT = 25 °C and 400 °C.
It needs to be emphasized that such a
peak occurs after hydrogenation of both
the “as-delivered” material and the weld
metal previously welded in a 7.5 vol.-%
hydrogen-argon atmosphere (Figure 11
a and b).
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A comparison of the TDS plots of the
electrochemically hydrogenated specimens (Figure 10) with those of other
specimens (Figures 8 and 9) reveals that
the amount of deuterium trapped inside
the material and desorbed at this temperature is higher than that of other
specimens.
Preliminary metallographic investigations reveal the existence of different
inhomogenities in the martensitic microstructure of these welds, and it is to be
expected that these defects act as various
trapping sites. For example, the largest
activation energy value obtained from
the respective desorption peak calculations stands at approximately 9 kJ/mole,
indicating that the respective trap belongs to a H-dislocation elastic stress field.
With respect to the various potential
trapping sites in supermartensitic stainless steel weld microstructures, the following considerations should be taken:
a) Desorption energy is defined as the
sum between the trap binding energy
and the activation energy for lattice
diffusion (Ed = Eb + Em). b) When the activation energy for lattice diffusion (Em)
is very large relative to the activation or
binding energy of a trap (Eb), a diffusion
controlled hydrogen evolution will be
measured.
As discussed above, the theoretical
and measured values of the activation
energy are numerically very close. This
in turn means that bulk diffusion is not
negligible and therefore that Lee and
Lee’s theory [7] does not apply in this
case. Lee and Lee’s theoretical approach,
which is based on hydrogen evolution
from trap sites only, is attractive because of its theoretical simplicity. However Lee and Lee’s model [7] does not
account for diffusion, correspondingly
it further does not account for the possibility of hydrogen released from one
trap becoming available to other traps.
Lee and Lee’s model is therefore unsuitable for the explanation of hydrogen
evolution and for the calculation of the
activation energy for the investigated
supermartensitic stainless steel weld
microstructures.
Conclusions
The investigation of the affects of different hydrogen exposures to welded
supermartensitic stainless steels has
lead to a better understanding of microstructural influences on the absorption/
52 (2010) 5
Specimen
Peak #
Temp. at
desorption
peak
(°C)
Maximal
desorption rate
(x1014 Hatoms/gr · s)
Half height
peak
width
(°C)
Calculated
activation
energy
(kJ)
I
30
11.30
50.78
4.2
II
92.34
3.39
37.50
5.1
III
130
2.80
25.78
6.5
7.90
51.23
4.56
MP
As-Received
As-Received
Ar + 0 % H2
Base Metal
Weld Metal
I
32.3
II
88.11
3.19
30.20
5.72
III
121
2.81
31.50
6.65
IV
143.85
2.63
28.43
6.68
I
30.94
8.74
50.31
4.54
II
82.34
3.40
28.59
5.70
III
110.94
2.94
27.97
6.62
IV
162.19
2.57
202.50
6.70
7.91
42.34
4.56
Ar + 7.5 % H2
Base Metal
Weld Metal
I
32.5
II
89.38
3.19
30.00
5.72
III
120
2.72
37.50
6.66
IV
143.75
2.57
28.59
6.68
I
34.84
9.33
47.03
4.57
II
92.5
3.56
36.56
5.73
III
117.5
3.24
27.19
6.63
IV
172.34
3.42
108.91
6.70
V
302.34
3.65
137.19
9.10
Table 4. Parameters of thermal desorption analysis from TIG welded SMSS, at a heating rate of
3 °C/min
Peak #
Temp. at
desorption peak
(°C)
Maximal desorption
rate
(·1014 D-atoms/g · s)
Half height
peak width
(°C)
As-Received charged with deuterium for 24 hours
I
77.5
112
38.44
II
155
167
172.50
Weld metal of TIG welded in Ar+7.5 % H2 charged with deuterium for 24 hours
I
80
128
40.63
II
152.34
195
148.13
Table 5. Parameters of thermal desorption analysis from SMSS electrochemically charged with
deuterium, for 24 hours at a heating rate of 3 °C/min
desorption behavior and the interaction
between hydrogen atoms and possible
trap sites. The following conclusions can
therefore be drawn:
1. The amount of absorbed hydrogen is
significantly dependent on the different modes of hydrogenation. Electrochemical charging can lead to much
higher levels of absorbed hydrogen
concentrations than charging in a
gaseous hydrogen atmosphere or during welding SMSS in a hydrogenated
shielding gas with up to 7.5 vol.-% H2.
2. The absorbed hydrogen concentration
significantly depends on the weld
microstructure. In the investigated
313
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mono-layer weld, lower concentrations are soluted in the HAZ as compared to the weld metal and to the
base material. This is to be attributed
to a respectively higher content of
retained and annealing austenite generated during production of the base
material and also to a higher amount
of retained austenite in the weld metal
itself. However, as the re-heating and
annealing processes produce respectively more austenite in the HAZ,
much higher hydrogen concentrations can be taken up in the HAZ
multilayer joints. The level of hydrogen in the weld microstructures is
thus also dependent on the joint design.
3. The influence of hydrogen rates in the
atmosphere on absorbed amounts appear only in atmospheres containing
7.5 % H2. Although there is some hydrogen loss in the period between
welding and final extraction procedures, if the atmosphere contains a
higher amount of hydrogen, more
hydrogen could be absorbed in the
material and be trapped there.
4. The changing of the lattice parameter
depends on the way in which hydrogen is introduced into the weld microstructures. For example, the lattice
parameters of the γ- and α-phase remain unchanged after hydrogenation
from a gaseous environment. However, the large amounts of hydrogen
introduced during electrochemical
charging procedures significantly reduce the α-phase reflections, while
the parameter of the γ-phase remains
unchanged due to the higher solubility of this phase.
5. The hydrogen evolution process was
found to be dependent on the hydrogenation mode. For specimens that
were subjected to hydrogen in the
atmosphere (as-received, TIG welded,
and gas phase hydrogenation) TDS
spectrums exhibited several evolution peaks. For specimens that were
subjected to electrochemical hydrogenation, hydrogen evolution was
observed in a large single peak that
covered a wide temperature range but
was located at lower temperatures.
This means that hydrogen is trapped
more deeply in the investigated SMSS
weld microstructures if introduced
during welding rather than during
other introductory procedures such
as cathodic charging.
314
Specimen
R.T.
650 °C
Peak #
Temp. at
desorption
peak
(°C)
Maximal
desorption rate
(x1014 Hatoms/gr · s)
Half height
peak
width
(°C)
Calculated
activation
energy
(kJ)
I
40.2
6.84
39.34
4.56
II
92.6
2.13
29.29
5.73
III
120.1
1.98
28.18
6.64
I
40
7.64
39.84
4.55
II
92.5
2.84
29.69
5.73
III
120
2.21
28.28
6.64
IV
285
2.99
358.28
7.59
Table 6. Parameters of thermal desorption analysis from SMSS, charged at 1 atm for 6.5 hours
at a heating rate of 3 °C/min
Abstract
Auswirkungen des Schützens mit verschiedenen WasserstoffArgon-Gemischen auf WIG-Schweißungen supermartensitischer
Stähle. Eine Anzahl verschiedener Effekte ergibt sich aus der Anwesenheit von Wasserstoff während des Schweißens hochlegierter
Stähle. Die Betriebsdauer von geschweißten Bauteilen ist außerdem
stark von der Anwesenheit von Wasserstoff im Umgebungsmedium
und der Anfälligkeit der verschiedenen Schweißnahtgefüge für eine
Degradation ihrer Eigenschaften durch Wasserstoff abhängig.
Als eine relative neue Werkstoffgeneration finden supermartensitische
hoch legierte Stähle (Supermartensitic Stainless Steels – SMSS)
zunehmend als Ersatz für teuere Legierungen insbesondere in der
Öl- und Gasindustrie Verwendung. Als Konsequenz ihres martensitischen Gefüges sind diese Legierungen anfällig für eine wasserstoffunterstützte Rissbildung (Hydrogen Assisted Cracking – HAC). Der
Widerstand von supermartensitischen Stählen gegen wasserstoffunterstützte Spannungsrisskorrosion (Hydrogen Assisted Stress
Corrosion Cracking – HASCC) unter Sauergasbedingungen wurde
vor allem für industrielle Einsatzzwecke extensiv untersucht. Solche
Studien vornehmlich an Grundwerkstoffen basieren überwiegend
auf Standard-Prüfverfahren. Dem gegenüber würde das grundsätzliche Verhalten von Wasserstoff in den Gefügen geschweißter supermartensitischer Stähle wenig untersucht. Die zentralen Gründe für
die diesem Beitrag zugrunde liegende Studie waren daher, die Effekte des Wasserstoffs auf das Gefüge von Wolfram Inert Gas (WIG)Schweißungen supermartensitischer Stähle und die entsprechenden
Wasserstoff-Trapping-Mechanismen zu untersuchen.
Die Wirkungen des Wasserstoffs auf die verschiedenen WIG-geschweißten Gefüge wurden mittels Röntgendiffraktometrie, Lichtmikroskopie und Rasterelektronenmikroskopie untersucht. Eine
Anzahl von Verfahren wurde außerdem angewendet, um den absorbierten Wasserstoff quantitativ zu bestimmen. Die Wechselwirkung
zwischen Wasserstoff mit den mikrostrukturellen Defekten und die
Charakteristika der Wasserstoffdesorption wurden mittels Ther52 (2010) 5
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mischer Desorptionsspektroskopie (TDS) und Trägergas-Heißextraktionen des Wasserstoffs (LECO Analyse) untersucht. Die Wirkung
des Gefüges auf die Absorption und Desorption von Wasserstoff
werden im Detail diskutiert.
Further investigation is required, which
will use alternative approaches and account for diffusion controlled hydrogen
transport in order to demonstrate which
trapping sites actually occur in the weld
microstructures.
References
1 Th. Boellinghaus: Hydrogen assisted cracking in supermartensitic stainless steels,
in: N. R. Moody, A. W. Thompson, R. Ricker,
G. Was, and R. Jones (Eds.) Hydrogen
Effects on Material Behavior and Corrosion
Deformation Interactions, TMS, Wyoming
(2003), pp. 1009-1018
2 D. Seeger; Th. Boellinghaus: Hydrogen permeation in supermartensitic stainless steel
weld microstructures, CORROSION 2004,
NACE International, Houston (2004),
paper 04142
3 E. Folkhard: Metallurgie der Schweißung
hochlegierter Stähle, Springer Verlag Wien
(1984), pp. 279-280 (in German)
4 Th. Boellinghaus, H. Hoffmeister: Hydrogen permeation in supermartensitic stainless steels dependent on heat treatment
and chemical composition, CORROSION
2000, NACE International, Houston (2000),
paper 00141
5 E. Tal-Gutelmacher, D. Eliezer, D. Eylon:
The effects of low fugacity hydrogen in
duplex- and beta-annealed Ti-6Al-4V alloy,
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(2004) pp. 230-236
6 E. Tal-Gutelmacher; N. Eliaz, D. Eliezer,
D. Zander, L. Jastrow, U. Köster: Absorption/desorption behavior of hydrogen and
deuterium in a Pd-Coated Zr-based amorphous alloy, Materials Science and Engineering A 358 (2003), pp. 219-225
7 S.-M. Lee; J.-Y. Lee: The trapping and transport phenomena of hydrogen in nickel,
in Metallurgical Transactions A, 17 (1986),
pp. 181-187
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8 P. Rozenak; D. Eliezer: Effects of metallurgical variables on hydrogen embrittlement
in AISI Type 316, 321 and 347 stainless
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MP
The Authors of This Contribution
Prof. Eliezer is a full Professor in the Department of Materials Engineering, Ben Gurion
University of the Negev (BGU), Israel, and
holds the Eric Samson Chair for Advanced
Materials and Processing at the BGU. Prof.
Eliezer is a fellow of the ASM (American
Society of Materials) (in recognition of distinguished contribution in the field of materials
science and materials engineering), and is a
Mercator Professor at DFG Clausthal.
Prof. Eliezer has held the position of
“Visiting Professor” and “Senior Associate” at
a number of Universities and Research Institutes across Europe, Asia, and USA, including
the: Air Force Wright Aeronautical Laboratories, National Research Council, USA;
NASA-Ames Research Centre, National
Research Council, USA; University of Illinois,
Department of Metallurgy and Mining, USA;
Federal Institute for Materials Research and
Testing (BAM), Germany, and the Research
Centre for Hydrogen Industrial Use and Storage (HYDROGENIUS), Japan. Prof. Eliezer
has published over 470 papers in Journals and
Conference Proceedings, and has edited numerous Scientific Books and Collective Volumes.
Yafit Nissim completed her B. Sc and
M. Sc at the Department of Materials Engineering at the Ben-Gurion University of the Negev,
Beer-Sheva, Israel. Her M.Sc Thesis “Hydrogen
Embrittlement of Welded Stainless Steel”
was completed in 2006. Yafit Nissim currently
works at Intel Israel.
Dr.-Ing. Thomas Kannengießer, born in
1971, studied Mechanical Engineering with the
background of Materials Engineering and Materials Testing at the University of Magdeburg.
From 1997 to 1999, he was a doctoral candidate at the BAM, Federal Institute for Material
Research and Testing, in Berlin and obtained
his doctoral degree with the subject “Investigations into the Formation of Welding-Specific
Stresses and Deformations at Variable Restraint
Conditions in the Component Weld Test”.
Since 2005 he is the head of the working
group “Component Testing” at BAM.
You will find the article and additional material by entering the document number MP110135
on our website at www.materialstesting.de
315