characteristics and effects of austenite resulting from
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Characteristics and effects of austenite resulting from
tempering of 13CrNiMo martensitic steel weld metals
P.D. Bilmesa,b,*, M. Solarib, C.L. Llorentea,c
aDepartamento de Mecanica, Facultad de Ingeniera, Universidad Nacional de La Plata, Calle 1 y 47, 1900 La Plata, ArgentinabConsejo Nacional de Investigaciones Cientficas y Tecnicas (CONICET), Av. Rivadavia 1917, 1033 Capital Federal, Argentina
cComision de Investigaciones Cientficas de la Provincia de Buenos Aires (CICPBA),
Calle 52, c/120 y 121, 1900 La Plata, Argentina
Received 31 July 2000; accepted 28 September 2000
Abstract
Low-carbon 13CrNiMo martensitic steels are remarkable for their high strength and high resistance to brittle
failure while retaining corrosion resistance together with weldability. These properties can be obtained when an
intercritical tempering is applied as heat treatment or postweld heat treatment (PWHT); promoting the
precipitation of finely distributed austenite that remains untransformed after cooling. The content and stability of
this austenite in the weld metal accounts for the high toughness even under subzero conditions. Transmission and
scanning electron microscopy (SEM), X-ray diffraction, and Mossbauer spectroscopy were used to study both the
austenite resulting from intercritical tempering of these soft martensitic stainless steel weld metals and the
austenitefracture interactions. To recognize the effect of the austenite content on impact toughness, single- and
two-stage tempering have been applied and evaluated through Charpy tests. The studies have shown the austenite
to be thermally stable, mainly due to its substructure, but not mechanically stable, indicating that the toughening
mechanism of the austenite particles is associated with transformation-induced plasticity (TRIP). D 2001 Elsevier
Science Inc. All rights reserved.
Keywords: Retained austenite; 13Cr-NiMo steels; Post weld heat treatments; Microstructure; Mechanical properties
1. Introduction
Soft martensitic stainless steels such as 13Cr
NiMo are widely used for hydraulic turbines, valve
bodies, pump bowls, compressor cones, impellers,
and high-pressure pipes in power generation, offshore
oil and gas, and petrochemical industries [1,2]. It is
known these steels perform well in applications
where corrosion and cavitation erosion resistance
are required. In addition, they have some resistance
to stress corrosion cracking (SCC) in CO2 and H2S
environments, high strength, and toughness even at
low temperatures or in thick cross-sections, and have
excellent weldability [36].
These alloys solidify to delta ferrite crystals [7].
The transformation of delta ferrite into austenite crys-
tals starts at around 1300C and ends, in the case ofequilibrium conditions, at around 1200C. With theactual cooling rates experienced during a welding
operation, small amounts of delta ferrite are super-
cooled during the delta ferrite) austenite transforma-
1044-5803/01/$ see front matterD 2001 Elsevier Science Inc. All rights reserved.
PII: S 1 0 4 4 - 5 8 0 3 ( 0 0 ) 0 0 0 9 9 - 1
* Corresponding author. Tel.: +54-221-423-6692; fax:
+54-221-425-9471.
E-mail address: [email protected] (P.D.
Bilmes).
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tion. Similarly, the austenite)martensite transforma-tion leads to a martensitic microstructure containing
small quantities of supercooled retained austenite.
Thus, after cooling has been completed, the structure
consists of soft martensite with small amounts of
supercooled delta ferrite and austenite.
When a subsequent intercritical tempering at
600C (slightly above the Ac1 temperature) is applied
as postweld heat treatment (PWHT), excellent tough-
ness properties are usually obtained. This temperingpromotes the softening of the martensite and the
precipitation of finely distributed austenite along the
martensite interlath boundaries and prior austenite
grain boundaries. After the tempering, this austenite
remains untransformed and it is known that these
particles account for the high toughness of this alloy.
However, the thermal and mechanical stability of the
austenite particles and the mechanism by which the
austenite enhances the toughness properties greatly is
not yet well understood. Several theories [8,9] have
attempted to explain how this happens in a compositemicrostructure of tempered lath martensite with small
particles of austenite phase densely distributed along
the lath boundaries and the prior austenite grain
boundaries. Among these theories, the crack blunting
model suggests that a crack being propagated through
the steel would blunt in the ductile stable austenite
(fcc). Although some evidence of this mechanism
was presented in a previous work [10], recent results
show that the transformation of austenite (fcc) into
martensite (bcc) happens during the fracture process
[11]. These new results are in agreement with themodels that are based on the transformation of
austenite into martensite by a localized transforma-
tion-induced plasticity (TRIP).
The present work shows the chemical and sub-
structural characteristics of the austenite particles,
present in 13Cr NiMo weld metals after intercritical
tempering, and their thermal and mechanical stability.
The effect of both single- and two-stage tempering on
the austenite content of the weld metal, and thus on
the impact toughness, is studied. Additionally, the
mechanism by which austenite particles improve thetoughness of these materials is recognized.
2. Experiments
Automatic gas metal arc welding was used to
prepare the weld metals. Multiple-pass welds were
performed on AISI 410 plates using a 13Cr NiMo
welding wire. The composition of the welding wire
and the welding parameters are shown in Tables 1
and 2. The applied PWHT were a single-stage tem-
pering (intercritical tempering at 600C/2 h/air) and atwo-stage tempering (first at 670C/2 h/air with a
previous solution annealing at 950C/1 h/air, second
at 600C/2 h/air). Chemical composition, microstruc-
tures, tensile properties, and Charpy V-notch impact
energy were determined in the as-welded and tem-pered conditions. The chemical composition of the
weld metal was measured by an optical emission
technique (except for C, N, O, and S that were
measured by combustion analysis). A Philips 515
scanning electron microscope (SEM) operated at 20
kV was used to observe the microstructures of the
PWHT weld metals and the fracture surfaces of the
Charpy V-notch impact specimens. Samples were
ground and electropolished, the electrolyte composi-
tion for the latter phase being: 62 ml HClO4, 700 ml
ethanol, 100 ml butyl cellusolve, and 137 ml H2O.The specimens were then etched in Vilellas solution.
The Charpy V-notch impact tests were performed at
20C and 77C.The volume fraction of the austenite in the tem-
pered conditions were measured by X-ray diffraction
from a Rietveld analysis. The carbon concentration in
this phase was evaluated by X-ray diffraction through
the lattice parameter of the austenite. X-ray diffraction
patterns were obtained at room temperature with a
Philips PW1710 diffractometer, furnished with a
diffracted beam graphite monochromator. Data werecollected using CuKa radiation in the range
10 2q 120 at a step interval of 0.02. A Rietveldanalysis was performed using the program DBWS-
9411 [12]. The sample displacement, the background
(modeled with a fifth-degree polynomial), the unit cell,
the preferred orientation, the pseudo-Voigt profile
parameters, and the scale factor of the different phases
present in the sample were refined independently but
not simultaneously. From the Rietveld analysis, a
relative weight fraction was assigned to the refined
Table 1
Nominal composition of welding wire (wt.%)
Grade Wire C Si Mn Cr Ni Mo
ASME SFA-5.9 class ER410NiMo 1.6 mm 0.03 < 0.95 0.65/0.9 12/13.5 5.5 0.5
Table 2
Welding parameters
Gas
Gas flow
(l/min)
Voltage
(V)
Current
(A) DC +
Arc speed
(mm/min)
Interpass
temperature
(C)
Ar 20 25 255 150 < 120
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Fig. 1. Ductile fracture surface of Charpy-V notch specimen: (a) Low magnification 1000, magnification bar is 10 mm, (b) High
magnification 6000, magnification bar is 10 mm.
Fig. 2. X-ray diffraction patterns of austenite: (a) Single-stage tempering, (b) Two-stage tempering.
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ical stability of austenite particles was also evaluated in
regions close to the fracture surfaces of Charpy speci-
mens in the plastic zone preceding the crack front. This
was performed by means of X-ray diffraction through a
Rietveld analysis. Microhardness was measured in
these zones by using a Shimadzu Vickers meter with
a load of 20 g and a loading time of 15 s.
3. Results and discussion
The chemical composition of the weld metal in
each of the different conditions is given in Table 3.The compositions were similar in all conditions and
in line with the nominal composition of the consum-
able electrode. The nitrogen and oxygen contents are
similar for both PWHT conditions. Tensile properties,
Charpy V-notch impact energy, hardness, and auste-
nite contents of the weld metals before and after
PWHT are shown in Table 4.
The yield and tensile strength were higher for the
as-welded martensitic condition. Two-stage tempering
promoted a greater softening than single-stage tem-
pering. The highest values of impact strength, or
toughness, were observed following the two-stage
tempering. This is the condition that contains the
greatest amount of austenite. As regards the fracture
mode, it is worth mentioning that all PWHT speci-mens tested both at room and subzero temperatures
displayed 100% of fibrous fractures with typical
Fig. 3. Difraction pattern with (y iobs) and (y i
cal) for Rietveld analysis. Two-stage tempering condition.
Fig. 4. Difraction pattern of M2X for two-stage tempering condition.
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dimples and without cleavage. Fig. 1a is representa-
tive of the fracture surfaces seen, this one from a
specimen tested at subzero temperature for single-
stage tempering. The typical dimple appearance asso-
ciated to ductile-dimple fracture with a high micro-
void density can be observed. This may be attributed
to the existence of the large number of internal
interfaces due to both nonmetallic inclusions and
austenite or transformed austenite particles, which
may act as void nucleation sites [14]. Fig. 1b shows
typical dimples with several small austenite or trans-formed austenite particles within them.
The X-ray diffraction patterns of the samples
corresponding to both single- and two-stage temper-
ing (Fig. 2a and b, respectively) were analyzed using
the DBWS-9411 program for Rietveld analysis. This
method was developed by Hugo Rietveld in 1969,
and today, it is one of the most powerful techniques
used for structural analysis, measuring reticular para-
meters, and other investigations about crystallogra-
phy. It consists of fitting step by step the experimental
intensities (yiobs) corresponding to all spectra, withthose (yi
cal) based on a specific crystalline structure
model, diffraction optical effects, instrument factors,
and other sample characteristics. The parameters
included in the model are refined to achieve the best
fitting of minimum squares of thousand of yi belong-
ing to the diffraction pattern. Fig. 3 shows the
diffraction pattern with (yiobs) and (yi
cal) correspond-
ing to two-stage tempered condition. Further details
of this applied technique can be found in Ref. [15].
According to the Rietveld analysis, the double-
tempered PWHT condition contained the higher
amount of retained austenite content (29.2 vol.%).As can be observed in Table 4, the higher austenite
contents are associated with the higher values of
impact toughness.
Carbonitrides of the type M2(C,N) were also
identified by means of X-ray diffraction in both single-
and two-stage tempered conditions. Fig. 4 shows the
diffraction pattern of the two-stage tempered condi-
tion, where the peaks corresponding to carbonitrides
have been noted. The presence of carbonitrides instead
of carbides of the type M7C3 or M23C6 is associated
with the high nitrogen content of the weld metalstogether with the presence of molybdenum. Usually
during tempering, the carbon is precipitated as car-
bide, and according to Irvine [16], precipitates such as
M3C, M2(C,N), M7C3, and M23C6 appear. According
to Pickering [17], the presence of nitrogen (and
molybdenum) promotes the formation of M2(C,N) at
the expense of M7C3 or M23C6.
Thus, both single- and two-stage tempering lead
to martensite decomposition together with the pre-
cipitation of a very thin austenite dispersion, since it
is known that the precipitation of austenite takes
Fig. 5. SEM micrograph of retained austenite between the
martensite laths. Single-stage tempering; 3000, magnifi-cation bar is 10 mm.
Fig. 6. SEM micrograph of retained austenite between the
martensite laths. Two-stage tempering; 3000, magnifica-
tion bar is 10 mm.
Fig. 7. Line scan analysis of carbon on an austenite particle;
20,000, magnification bar is 0.2 mm.
Table 5
Average chemical composition of austenite particles as
measured by TEM/EDS (wt.%)
Cr Ni Mn Si Mo Fe
13.7 8.3 1.9 0.8 1.3 balanced
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place at tempering temperatures slightly higher than
Ac1 (600C). This austenite is shown in Figs. 5 and 6,
for single- and two-stage tempered conditions respec-
tively, arrayed like platelets among martensite laths. It
remains stable and does not later transform intomartensite during the cooling after the tempering.
The line scan analysis of carbon on an austenite
particle in the single-stage tempered condition, deter-
mined by EPMA, is shown in Fig. 7. This indicates the
presence of a considerable amount of carbon. The
actual content level has been determined by X-
ray diffraction from lattice parameter measure-
ments. The austenite lattice parameter measured was
aog= 0.35901 nm, from which the carbon content has
been calculated using Eq. (1) from Refs. [18,19]:
aog 3:572 0:033 wt:% C: 1
According to this equation, the carbon concentration
is calculated to be 0.548 wt.%.
The approximate composition of the austenite
phase as evaluated by EDS analysis using TEM is
presented in Table 5. This indicates that the austenite
particles are enriched in solute elements such as Ni,
Mn, and Mo, when compared with the nominal
composition of the welding wire (Table 1). This
solute enrichment in the austenite particles will leadto a decrease in the Ms temperature. The Ms tempera-
ture at which the precipitated austenite would begin
to transform spontaneously into martensite has been
calculated in terms of the austenite composition by
Eq. (2), which is specific to soft martensitic stainless
steels (from Ref. [7]).
MsC 492 125 wt:% C 65:5
wt:% Mn 10 wt:% Cr
29 wt:% Ni 2
The Ms temperature thus calculated is 78C.
The microstructure after single-stage tempering
observed by TEM is shown in Figs. 8 and 9. The
fine austenitic platelets are the darkest platelet-like
regions with an increased dislocation density. Other
investigators have observed a similar substructure
with localized carbide particles and orientation rela-
tionships between matrix and austenite, such as the
KurdjumovSachs relationship, in similar soft mar-
tensitic stainless steels [20,21]. All these character-istics suggest that the austenite could have been
formed by a shear process [22].
The Mossbauer spectra of a specimen cooled at
different temperatures down to 196C and held atthis temperature for 20 h are shown in Fig. 10. All
spectra reveal six lines belonging to the different
ferromagnetic phases and a central line, corresponding
to a paramagnetic phase. This central peak corresponds
to the austenite phase. No changes in the austenite
content were noted with decreasing temperature. This
indicates that the austenite particles did not transforminto martensite and retained their thermal stability even
after cooling to 196C.The Ms determined for the enriched retained
austenite was 78C. This value is not in agreementwith the results obtained by subzero treatments.
Although other equations were used to calculate Ms,
the results were similar to that obtained using Eq. (1).
This indicates that the solute enrichment of the
austenite particles could be a factor contributing to
the stability of the particles, but it is not the only
factor accounting for the austenite thermal stability.Accordingly, it is suggested that the stability of the
austenite has substructural as well as chemical ori-
gins. Therefore, there is a possibility that the
observed substructure in the austenite could increase
its stability against the transformation into martensite
on cooling. The stability of the austenite particles is
probably due to increased difficulty in propagatingFig. 8. TEM micrograph of retained austenite between the
martensite laths; 43,000, magnification bar is 1 mm.
Fig. 9. High dislocation density within the retained
austenite particles; 75000, magnification bar is 0.5 mm.
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the shear of the martensitic transformation via the
increased barriers afforded by the substructure in the
austenite particles.
In order to test the mechanical stability of the
austenite particles, some part of the material was
cold rolled at room temperature to reductions inthickness of 20%, 40%, 60%, and 80%. The
mechanical stability of the austenite is defined as
the susceptibility of austenite to transform into
martensite, as induced by plasticity. An austenite of
higher mechanical stability needs more strain to
transform into martensite than that of lower mechan-
ical stability. The changes in the austenite content
associated with the austenite (fcc) to martensite (bcc)
transformation were assessed by Mossbauer spectro-
scopy and are shown in Fig. 11. The spectra show
that with only 20% reduction in thickness some ofthe austenite has been transformed into martensite.
Above 60% reduction in thickness the central lines
corresponding to austenite are absent, so by this
stage practically all the austenite has been trans-
formed mechanically into martensite.
After Charpy V-notch impact testing, regions
close to and far from the fracture surfaces of the
Charpy V-notch specimens in the single-stage tem-
pered condition, were analyzed by X-ray diffraction
to study the interaction between the fracture front
and the austenite particles. This interaction could
result in the transformation of the austenite particlesinto martensite. The X-ray diffraction patterns of
these regions corresponding to single-stage tempered
condition are shown in Fig. 12a and b. The austenite
diffraction peaks were observed in regions far from
the fracture but not in regions close to the fracture
surface. This indicates that austenite particles were
transformed into martensite in regions close to the
fracture. In order to view effectively the microstruc-
ture immediately adjacent to the fracture, the fracture
surface of the Charpy specimen was electrolytically
covered with metallic chromium. An intense strainfield was observed close to the fracture by SEM
(Fig. 13) together with microvoids around trans-
formed austenite particles (Fig. 14). The average
microhardness in this region adjacent to the fracture
Fig. 10. Mossbauer spectra before and after the subzero
treatment.
Fig. 11. Mossbauer spectra before and after cold rolling at
room temperature.
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Fig. 12. X-ray diffraction patterns performed close to and far from the fracture surfaces of the Charpy-V notch specimens for
single-stage tempering: (a) Region far from the fracture surface of Charpy-V notch specimen, (b) Region immediately close to
the fracture surface of Charpy-V notch specimen.
Fig. 13. Intense strain field closer to the fracture surface of
Charpy-V notch specimen; 3000, magnification bar is10 mm.
Fig. 14. Microvoids around transformed austenite particles
close to the fracture of a Charpy specimen; 15,000,
magnification bar is 1 mm.
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surface was 320 HV, which is near the value of the
as-welded condition where the microstructure ismainly martensite. However, far from the fracture,
the average microhardness was 275 HV, in accor-
dance with the value of the single-stage tempered
condition (Table 4).
All these results suggest that the austenite retained
particles suffered the martensitic transformation dur-
ing the progression of a crack front and they could
have acted as energy absorbers. In the event of a
propagating crack passing into or near metastable
regions, the concentrated strain field at the crack tip
enables austenite particles to transform into stable,but less dense, martensite. This transformation,
mechanically induced in the plastic zone, absorbs
additional energy, thus effectively enhancing the
toughness. The associated volumetric expansion of
this transformation tends to close the crack and
relieve stresses at its tip, absorbing strain energy
during the fracture that might otherwise have gone
through the crack extension. This kind of transforma-
tion mechanism is recognized to be primarily respon-
sible for the beneficial toughening effect of a
metastable phase within the microstructure [23].Fig. 6 shows the microstructure of the material in
the double tempered condition (first at 670C/2 h/air,second at 600C/8 h/air). A structural refinement can
be noted, meaning a higher austenite content together
with a more uniform distribution of this phase. The
mechanism by which austenite precipitation increases
with double tempering is thought to be associated
both with the instability of the austenite particles
during cooling following tempering at 670C and tolonger times of tempering at 600C. The austenite thatwas produced at 670C was not stable enough so it
became partially transformed into martensite during
cooling in air. Fig. 15 shows the diffraction pattern of
a X-ray diffraction analysis performed on a specimen
after the first tempering at 670C. The principal peak
of austenite is barely apparent. In this condition the
microstructure shows a second precipitated phase
with similar morphology like the austenite particles
(Fig. 16). In fact, this precipitated phase is mainly
Fig. 15. Diffraction patterns of a X ray diffraction analysis performed on a specimen after the first tempering at 670 C/2h/air.
Peaks of austenite barely appear.
Fig. 16. Precipitated martensite with similar morphology to
the austenite particles, resulting from the martensitic
transformation of the austenite particles which are not stable
enough, after the first tempering at 670C/2h/air; 3000,
magnification bar is 10 mm.
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martensite, resulting from the martensitic transforma-
tion of the only partially stable austenite particles.
Thus, after the first tempering, the microstructure was
composed of tempered martensite, some retained
austenite, and some fresh martensite. During thesecond tempering at 600C, new stable austenite
particles platelets precipitate through new interfaces
(fresh martensite/austenite) and the fresh martensite is
decomposed into tempered martensite. A similar
mechanism has been proposed by other investigators
in these steels [20]. On the other hand, longer temper-
ing times, 8 h, at 600C (four times longer than in
single-stage tempering) promoted a very high content
of retained austenite, more than in single-stage tem-
pered condition. Hence, as result of this two-stagetempering, an increase in the number of austenite
platelets after double tempering was produced (29.2
vol.% according to Table 4) and with a more uniform
distribution. A scheme of this structural refinement
due to two-stage tempering is shown in Fig. 17.
4. Conclusions
1. The studies have shown that the microstructure
of 13% Cr NiMo weld metals, after single- andtwo-stage tempering, consists of tempered martensite
and retained austenite with an acicular or lath-like
morphology, very thinly spread along the martensite
laths. These austenite particles are enriched in C, Ni,
and Mn, and have a high dislocation density. The
lath-like austenite morphology and the substructure
of the austenite phase, suggest that the austenite
par ticles could be for med by a shear process
between the laths of martensite that have been
enriched in stabilizing elements during the intercri-
tical tempering.2. According to the results from treatments at
different subzero temperatures, the austenite particles
were thermally stable. However, they could be
transformed mechanically into martensite by cold
rolling. Mechanical transformation was also observed
close to the crack front of the Charpy V-notch
specimens where localized plastic deformation took
place. This indicated the austenite had been trans-
formed by TRIP.
3. The solute enrichment of the austenite parti-
cles may be a factor contributing to the low Ms ofthe particles, but it does not fully account for the
high austenite thermal stability. The presence of a
high dislocation density within the austenite parti-
cles suggests that the stability of these particles may
have substructural as well as chemical origins.
Some stability of the austenite particles may be
probably due to increased difficulty in propagating
the shear of the martensitic transformation via
increased barriers afforded by the substructure in
the austenite particles.
4. The yield and tensile strength values werehighest for the as-welded condition. By comparison,
the greatest softening was induced by the PWHT that
involved a two-stage temper treatment. This heat
treatment also produced the largest amount of
retained austenite. Regarding the mechanical proper-
ties, a good balance was achieved between maximum
toughness, high ductility, and high strength.
Fig. 17. Scheme of the structural refinement due to two-stage
tempering: (a) 950C/1h/air, Martensite lath (MI), (b) 950C/
1h/air + 670C/during heating. Tm: tempered martensite
g: austenite, (c) 950C/1h/air + 670C/2h/air (after cooling).
Tm: tempered martensite; MI: Martensite lath, g: austenite,
(d) 950C/1h/air + 670C/2h/air + 600C/2h/air. Tm:
tempered martensite; g: austenite.
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