metals

Article
Microstructural Heterogeneity and Mechanical Properties of a
Welded Joint of an Austenitic Stainless Steel
Jairo Alberto Muñoz 1, *, Egor Dolgach 2 , Vanina Tartalini 3 , Pablo Risso 3 , Martina Avalos 3 , Raúl Bolmaro 3
and José María Cabrera 1,4

1 Department of Materials Science and Engineering EEBE, Universidad Politécnica de Catalunya, c/Eduard
Maristany 10–14, 08019 Barcelona, Spain
2 Laboratory of Hybrid Nanostructured Materials, National University of Science and Technology “MISIS”,
Moscow 119049, Russia
3 Instituto de Física Rosario, Consejo Nacional de Investigaciones Científicas y Técnicas-CONICET,
Universidad Nacional de Rosario, Ocampo y Esmeralda, Rosario 2000, Argentina
4 Fundació CIM-UPC, c/Llorens i Artigas 12, 08028 Barcelona, Spain
* Correspondence: jairo.alberto.munoz@upc.edu

Abstract: This research presents the microstructural and mechanical evolution throughout the welded
seam of an austenitic stainless steel (ASS) tube. It was found that the main hardness decrement
occurred in the fusion zone (FZ), followed by the heat-affected zone (HAZ) and the base material (BM).
Optical microscopy indicated a dendritic structure in FZ and heterogeneous austenitic grain size from
the HAZ towards the BM, ranging from 100 µm to 10 µm. The welding process generated an intense
texture around the FZ and the HAZ, while the BM still showed an extrusion-like texture. In terms of
mechanical behavior, the largest austenite grain size in the FZ led to the lowest strength and ductility
of all zones due to the earliest strain localization manifested by heterogeneous strain distribution.
However, the strain localization in all zones appeared after 0.4 true strain, indicating an overall good
ductility of the seam. These high values were related to two microstructure characteristics: (1) the
10% δ-ferrite after solidification in the FZ favored by the Creq /Nieq = 1.67 relationship that delayed
the crack propagation along the austenite grains and (2) the heterogeneous microstructure made up
of soft austenite and hard martensite in the HAZ and BM producing multiple strain concentrations.
Citation: Muñoz, J.A.; Dolgach, E.;
Kernel Average Misorientation (KAM) maps obtained by Electron Back-Scattering Diffraction (EBSD)
Tartalini, V.; Risso, P.; Avalos, M.;
allowed observing higher internal misorientations in the FZ than in the HAZ due to interconnected
Bolmaro, R.; Cabrera, J.M.
Microstructural Heterogeneity and
walls between the δ-ferrite grains. However, the largest KAM values were observed in the BM
Mechanical Properties of a Welded between γ-austenite and the deformation-induced α’-martensite phases. X-ray diffraction revealed
Joint of an Austenitic Stainless Steel. that the residual stresses in the cross-section of the welded seam were compression-type and then
Metals 2023, 13, 245. https://doi.org/ switched to tension-type in the outer surface.
10.3390/met13020245
Keywords: austenitic stainless steel; microstructure; strength; welding; heterogeneity
Academic Editor: C. Issac Garcia

Received: 30 December 2022

Revised: 24 January 2023
Accepted: 26 January 2023 1. Introduction
Published: 27 January 2023
Welding as a technological manufacturing process that requires knowledge of the
properties across the welded seam to fulfill the load applications, corrosion resistance,
fatigue strength, and safety requirements, i.e., the joint quality. It is so that welding process
Copyright: © 2023 by the authors.
investigations represent a comprehensive research field in metallurgy and materials science,
Licensee MDPI, Basel, Switzerland. where it is sought a better understanding of different hardening and softening phenomena,
This article is an open access article as well as the relationship between processing parameters and microstructure changes that
distributed under the terms and influence the welded joint integrity. Therefore, chemical compositions of the elements to
conditions of the Creative Commons joint, type of welding process, protecting atmosphere, heat input, and filler material are
Attribution (CC BY) license (https:// vital parameters defining the quality of the welded seam [1–3].
creativecommons.org/licenses/by/ Within the significant number of welding processes, Gas Tungsten Arc Welding
4.0/). (GTAW), also known as Tungsten Inert Gas (TIG), stands out due to its ability to op-

Metals 2023, 13, 245. https://doi.org/10.3390/met13020245 https://www.mdpi.com/journal/metals

Metals 2023, 13, 245 2 of 18

erate with and without filler material, allowing to obtain sound quality and free-of-defects
weld seams of similar and dissimilar materials [4]. These advantages make the GTAW
process ideal to be used in several industrial applications, such as applying an overlay, fill
pass welding, and welding by root joints [4,5]. For these reasons, the GTWA process has
been used to manufacture seam tubes, repairs, and joints between tubes to increase their
length [6].
When it comes to high-quality and performance tubes, steels offer a wide diversity
of properties that fit many applications. For example, stainless steels include a group of
alloys covering an ample and diverse range of properties, starting with corrosion resistance,
formability capacity, and thermal stability at high and cryogenic temperatures. Additionally,
they are inert materials, and due to their low carbon content, most stainless steels show good
welding behavior [7–9]. These distinct properties are highly related to alloying elements
such as C, Cr, and Ni that define the steel phases and the solidification mechanism [10,11].
Stainless steel weld seam tubes can be found in many industrial types of equipment.
Some examples are the chemical industry, petroleum and gas transport, paper factories,
the electric industry, medical implements, the food industry, construction and buildings,
aerospace applications, textile industries, and nuclear and power plants [12–14]. Thus, the
study, comprehension, and evaluation of the weld seam tubes’ properties under different
environments and working conditions are essential to ensure optimal performance.
Some authors have studied the effect of different processing variables and how these
variables affect the intergranular and pitting corrosion resistance in stainless steel, con-
sidering different temperature ranges and measurement techniques [15,16]. For example,
Kain et al. [17] studied the stress corrosion cracking phenomenon in stainless steels, evalu-
ating the composition, stacking fault energy, microstructure, and environmental conditions.
In this study, they concluded that a high fraction of low-energy grain boundaries was
helpful to counteract the stress corrosion cracking sensibilization. On the other hand,
Böhner et al. [18] evaluated the effect of the martensitic transformation on the steel fatigue
life applying monotonic and cyclic loads. Their principal findings were that under cyclic
loads, the critical nuclei for the martensitic transformation were independent of ultrafine-
grained microstructure. However, the opposite happened with monotonic loads, where
the ultrafine-grained microstructure was more favorable for the martensitic transformation
than its coarse-grained counterpart.
Martensitic transformation in stainless steels is of great importance because the steel
can combine hard and soft phases such as martensite and austenite that help to improve the
steel fatigue life delaying the crack propagation due to a local hardening effect [19]. This is
why studying the phase transformations in the welding process is a key parameter defining
the joint properties. For example, Lima et al. [20] demonstrated grain growth during the
temperature rise in the Heat Affected Zone (HAZ) and the Cr carbides nucleation that
diminished the joint mechanical strength. Moreover, Muthupandi et al. [21] proved that
the chemical composition of duplex stainless steel has a stronger influence on the ferrite–
austenite relation than the cooling rate during the welding process. Several authors have
pointed out that after solidification, the ferrite solidification in the weld seam of austenitic
stainless steel alloyed with CrMnNi helps to improve the hot cracking resistance [22,23].
Welding processes involve considerable microstructural modifications around the
weld seam. Thus, different zones are identified. For instance, the Fuzion Zone (FZ), where
the highest temperatures are reached because the fusion of the materials is necessary, giving
rise to a dendritic microstructure. Then, the HAZ, where phase transformations and grain
growth can occur. Therefore, studying the microstructure evolution, mechanical properties,
and the mechanism behind the welding process of stainless steel is a topic of interest
that guarantees the integrity of elements to avoid failure. In this context, it is crucial to
determine in detail the properties of each region from the FZ to the Base Material (BM)
passing thorough the HAZ.
This investigation addresses the microstructural and mechanical heterogeneity of a
weld seam obtained by the GTAW in austenitic stainless steel (ASS) 304L. To apport a
 properties, and the mechanism behind the welding process of stainless steel is a topic of
interest that guarantees the integrity of elements to avoid failure. In this context, it is cru-
cial to determine in detail the properties of each region from the FZ to the Base Material
(BM) passing thorough the HAZ.
Metals 2023, 13, 245 This investigation addresses the microstructural and mechanical heterogeneity 3 of
of18a
weld seam obtained by the GTAW in austenitic stainless steel (ASS) 304L. To apport a
better understating of the steel properties subject to a welding process, the microstructure
heterogeneity
better between
understating thesteel
of the FZ and the BMsubject
properties was thoroughly studied
to a welding by Optical
process, Microscopy
the microstructure
heterogeneity betweenmicroscopy
(OM) and electronic the FZ and using
the BM was thoroughly
Electron studiedDiffraction
Backscattering by Optical(EBSD).
MicroscopyFur-
(OM) and electronic
thermore, hardness microscopy
profiles andusing Electron Backscattering
micro-tensile Diffraction
tests for each zone (EBSD). Further-
were evaluated regard-
more,
ing the hardness profiles
weld joint and micro-tensile
mechanical tests for
behavior. Thus, theeach zonepoints
critical were and
evaluated regarding sof-
the responsible the
weld joint mechanical behavior. Thus, the critical points
tening and hardening mechanisms across the weld seam were established.and the responsible softening and
hardening mechanisms across the weld seam were established.
2. Materials and Methods
2. Materials and Methods
2.1. As-Received Material
2.1. As-Received Material
The as-received material was an extruded tube of 100 mm exterior diameter and 4
The as-received material was an extruded tube of 100 mm exterior diameter and 4 mm
mm in thickness. The BM steel corresponds with a 304L ASS with the chemical composi-
in thickness. The BM steel corresponds with a 304L ASS with the chemical composition
tion indicated in Table 1. The tube was cut and welded around its perimeter, as shown in
indicated in Table 1. The tube was cut and welded around its perimeter, as shown in
Figure 1a. A small cross-section of the weld seam in the radial direction (RD)-extrusion
Figure 1a. A small cross-section of the weld seam in the radial direction (RD)-extrusion
direction (ED) plane was extracted to analyze the microstructure and the different zones
direction (ED) plane was extracted to analyze the microstructure and the different zones
generated by the heat transfer, as indicated in Figure 1b and Figure 1c, respectively. The
generated by the heat transfer, as indicated in Figure 1b,c, respectively. The GTWA process
GTWA
in a girth process in a girthwith
configuration configuration
a current with
of 250 a current of 250
A, voltage of A,
27 voltage of 27 V, and non-
V, and non-consumable
consumable 2.4 mm diameter electrode was employed to
2.4 mm diameter electrode was employed to weld the pieces. Argon flowing weld the pieces. Argon
at aflowing
rate of
15 L/min was used as a protector atmosphere, and the Filler Material (FM) with a(FM)
at a rate of 15 l/min was used as a protector atmosphere, and the Filler Material with
chemical
a chemical composition indicated in Table 1 (we did not apply weld
composition indicated in Table 1 (we did not apply weld root protection to purge the root protection to
purge the
oxygen oxygen
inside inside the
the tube). Thetube).
main The main differences
differences between the between the baseand
base material material and
the filler
the filler material are concentrated in the Mn content, which is considerably
material are concentrated in the Mn content, which is considerably higher in the filler higher in the
filler
material.material.

Table1.1.Chemical
Table Chemicalcomposition
compositionforfor
thethe studied
studied steel
steel base
base material
material (BM)
(BM) andand
fillerfiller material
material (FM)(FM) (wt.
(wt. %).
%).
Zone C Cr Mn Ni Si P S Fe
Zone C Cr Mn Ni Si P S Fe
BM 0.01 17.1 0.09 9.1 1.1 0.03 0.01 Bal.
FM 0.06
BM
18.8
0.01
1.6
17.1 9.8 0.09 9.1
1.3
1.1
0.04
0.03 0.03 0.01 Bal.
Bal.
FM 0.06 18.8 1.6 9.8 1.3 0.04 0.03 Bal.

Figure1.1. (a)
Figure (a) welded
welded tube,
tube, (b)
(b)welded
weldedseam
seamsection,
section,and
and(c)
(c)welded
weldedseam
seammicrostructure.
microstructure.

2.2. Microstructure Characterization

The microstructure of a representative area of the welded joint and the BM was
characterized by cutting the tube along its longitudinal direction (LD) to observe the weld
seam, and BM cross sections, as indicated in Figure 1. Microstructural changes in the FZ,
HAZ, and BM were observed through OM using an optical microscope Zeiss AX10 (ZEISS
Microscopy, Oberkochen, Germany), as shown in Figure 2. The sample surfaces were
 with the addition of one more polishing step using 0.03 µm particle size colloidal silica
that does not crystalize. This procedure was carried out in a vibratory polishing machine
for 3 h and subsequently cleaned in ultrasonic for 1 h before taking the sample inside the
microscope chamber. The EBSD scans were obtained in different zones between the weld
seam and BM to reveal the microstructure heterogeneity, as displayed in Figure 2. For the
Metals 2023, 13, 245 texture and microstructure analysis, the EBSD scans were obtained with 1 µm and 0.1 4 ofµm
18
step sizes, respectively. It is worth mentioning that the indexation rate in all the scans was
over 90%, so the non-indexed points were corrected by the grain dilation method, and
grains with
prepared less than 2 pixels
by mechanical were
grinding excluded
following from any
standard statistical analysis.
metallographic The data
procedures. As awere
first
processed
step, by TSL-OIM
the surfaces 7.1 andwith
were polished MTEX toolbox software
SiC sandpapers (version 5.7.0,ranging
with granulometry Ralf Hielscher,
between
Freiberg,
400 Germany)
and 2500. Then, a[25]. Grain boundaries
fine polishing were divided
with diamond into of
suspensions two families:
9 µm, 6 µm, high
3 µm,angle
and
1grain
µm was applied.(HAGB)
boundaries Then, the sample
for those surface was with
boundaries electropolished with 10%
misorientations aqueous
higher than 15°oxalic
and
acid
low at room
angle temperature
grain boundaries for(LAGB),
5 s to reveal
alsothe microstructure
referred and to with
to as subgrains, minimize any phase
misorientations
transformation
between 3° andinduced by local
15°. Grain deformations.
size and Finally,Misorientation
Kernel Average ImageJ software was used
(KAM) estimate
profiles were
the average
plotted grain size
following the from optical micrographics
methodology proposed by by derivation
Muñoz of theFor
et al. [26]. equivalent
the sakediameter
of statis-
of each
tics, grain
scan [24].
areas of 500 µm × 500 µm were considered for the texture analysis.

Figure2.2. Analyzed
Figure Analyzed zones
zones on
on the
the weld
weld seam
seam with
withindication
indication of
ofthe
theareas
areaswere
weremini-tensile
mini-tensilesamples
samples
were extracted (all the tensile samples were cut with the same orientation; the tensile test drawings
were extracted (all the tensile samples were cut with the same orientation; the tensile test drawings
are scaled 1:2).
are scaled 1:2).

For a more detailed microstructure, texture, and phases transformations characteri-

zation on the different zones, a TSL-OIM EDAX EBSD system mounted on an FEG SEM
Quanta 200 scanning electron microscope (SEM, manufactured by ThermoFisher scientific,
Waltham, MA, USA) was used. The sample preparation was similar to the OM, but with
the addition of one more polishing step using 0.03 µm particle size colloidal silica that
does not crystalize. This procedure was carried out in a vibratory polishing machine for
3 h and subsequently cleaned in ultrasonic for 1 h before taking the sample inside the
microscope chamber. The EBSD scans were obtained in different zones between the weld
seam and BM to reveal the microstructure heterogeneity, as displayed in Figure 2. For the
texture and microstructure analysis, the EBSD scans were obtained with 1 µm and 0.1 µm
step sizes, respectively. It is worth mentioning that the indexation rate in all the scans
was over 90%, so the non-indexed points were corrected by the grain dilation method,
and grains with less than 2 pixels were excluded from any statistical analysis. The data
were processed by TSL-OIM 7.1 and MTEX toolbox software (version 5.7.0, Ralf Hielscher,
Freiberg, Germany) [25]. Grain boundaries were divided into two families: high angle
grain boundaries (HAGB) for those boundaries with misorientations higher than 15◦ and
low angle grain boundaries (LAGB), also referred to as subgrains, with misorientations
between 3◦ and 15◦ . Grain size and Kernel Average Misorientation (KAM) profiles were
plotted following the methodology proposed by Muñoz et al. [26]. For the sake of statistics,
scan areas of 500 µm × 500 µm were considered for the texture analysis.
In addition to the OM and EBSD studies, X-ray diffraction analyses were carried out
on the FZ and the BM zones. The diffractograms for each zone were obtained by means
of a Philips X pert Pro MPD diffractometer with Cu-Kα radiation at 40 kV and 15 mA, a
polycapillary fiber optic concentrator device (called X-ray Lens by Panalytical, Malvern
Panalytical Ltd, Malvern, UK), and Soller slits on the outgoing X-ray beam reaching a Xe
gas detector. The whole setup has the advantage of a high intensity parallel beam with
Metals 2023, 13, 245 5 of 18

superb averaging capabilities. Residual strains were measured, and stresses were calculated
by using the proper polycrystalline isotropic elastic constants for 304 stainless steel by
the sin2 ψ method. The method allows the measurement of stresses without knowing the
precise value of dhkl o (interplanar distance) for different family planes. That is particularly
useful for studying materials with low anisotropy as FCC homogeneous metals and alloys.
The intensities in the 2θ range between 38◦ –118◦ were measured with a scan step size of
0.02◦ . For the volumetric fraction calculation of each phase, it was assumed proportional to
the integrated intensity of all its peaks, as indicated by the following equation [27]:
j
n
1
∑nj=1 Iα0
Vα0 = j j j
(1)
1 n
n ∑ j=1 Iα0 + 1
n ∑nj=1 Iγ /Rγ

where n represents the number of integrated peaks for each phase; I is the integrated
intensity; R refers to the scattering factor. All the diffractograms were analyzed using Maud
software [28], which allows for determining the dislocation density (ρ) considering the
peak broadening according to the following equation [29]:
√
2 3(ε2 )1/2
ρ= (2)
b·d
where ε is the lattice strain; b the Burgers vector; d is the crystal size.

2.3. Mechanical Characterization

Uniaxial tensile tests and hardness measurements evaluated the mechanical response
for all the zones. Vickers hardness profiles were built inside the FZ, across the fusion line
between the FZ and the HAZ and on the BM for better mechanical characterization. All
the hardness measurements used a load of 1.97 N and a dwell time of 15 s. To avoid the
strain field influence from the previous hardness mark, every new hardness measurement
was taken at a distance of two times the size of the last mark. For the uniaxial tensile
tests at room temperature, two scaled bone shape samples with gauge dimensions of
4 mm × 1.3 mm × 2 mm were extracted from each zone (see Figure 2) using electrode
discharge machining (EDM). All the tensile samples were tested in a universal Instron
machine with a constant strain rate of 1 × 10−3 s−1 . During the tensile tests, the strain and
the deformation maps were recorded using the digital image correlation (DIC) method. In
this regard, a speckle pattern of black and white spots was painted on the calibrated length
of each sample, and pictures were taken every 2 s. The post-processing was conducted
with the free software Ncorr (version v1.2.2, Georgia Institute of Technology, Atlanta, GA,
USA, http://www.ncorr.com/index.php) [30].

3. Results and Discussion

3.1. Microstructure Evolution
Figure 3a shows the weld seam transversal section after chemical etching, where the
main characteristics of the FZ and the HAZ microstructures can be identified. Figure 3b
highlights the microstructure profile inside the FZ following the dashed red arrow in
Figure 3a. In this profile, distinct microstructural characteristics of the solidified material
are noticed. For example, the microstructure in the FZ is fully dendritic in all the sections.
Second, the dendritic structure solidified following different behaviors depending on the
zone. Larger and narrower dendrites were observed in the middle zone of the FZ than close
to the fusion line, where a more equiaxed cellular was observed, as corroborated by the
black square insets in Figure 3b. Some authors have found dendritic microstructures in the
FZ with cellular and columnar morphologies, either for the δ-ferrite (bcc crystal structure)
or for the γ-austenite (fcc crystal structure) in ASS after the welding process [31,32]. The
coexistence of various dendritic morphologies in the FZ has been associated with local
chemical fluctuations and cooling rate gradients across the FZ [22].
Metals 2023, 13,2023,
Metals 245 13, 245 7 of 19 6 of 18

Figure 3. (a) Weld seam cross section indicating the FZ, HAZ, and BM, (b) microstructural pano-
Figure (a) Weld
3. across
ramic theseam
FZ, (c)cross section indicating
microstructural panoramicthebetween
FZ, HAZ,the and
HAZBM,
and(b)
themicrostructural panoramic
BM, and (d) austenitic
across thesize
grain FZ,evolution
(c) microstructural panoramic
across the interphase between
of the FZ andthe
the HAZ and the
HAZ until BM, and (d) austenitic grain
the BM.
size evolution across the interphase of the FZ and the HAZ until the BM.

Regarding the HAZ and BM microstructure, Figure 3c confirms significant grain size
variations from the FZ-HAZ interphase until zones beyond 1 mm far from the fusion
line. As expected, heat transfer affects areas closer to the FZ more than the furthest ones.
This gave rise to austenitic grain sizes next to the fusion line of 100 µm. Figure 3d plots
three grain size profiles following the green dashed arrows indicated in Figure 3a to
corroborate the grain size variations. These profiles suggest the existence of grain sizes
larger than 100 µm around the FZ-HAZ neighborhoods, which continuously decrease as
one moves away from the fusion line reaching average values of 10 µm in the BM. It is
worth mentioning that grain size values calculated from OM mainly correspond to the
austenite phase. However, the grain size and grain boundaries from other phases, such
as δ-ferrite and α’-martensite (bcc crystal structures) coming from the solidification and
Metals 2023, 13, 245 7 of 18

forming processes, respectively, are not easy to reveal and identify by OM. Even so, it is
evident from Figures 3c and 3d that the temperature gradient creates an HAZ that extends
beyond 1 mm from each side of the weld seam.
Figure 4 shows SEM images for all the analyzed zones. Figure 4a confirms the so-
lidification of the primary δ-ferrite together with the austenite grains. In the fusion line,
Figure 4b allows observing several phases with different morphologies. Austenite is mainly
observed on the FZ side, and new microstructural characteristics are noticed on the HAZ.
These new characteristics are related to the presence of α’-martensite induced by the tube-
forming process. Thus, in the BM zone, two phases can be distinguished following a
Metals 2023, 13, 245 8 of
well-defined orientation, as indicated in Figure 4c. Furthermore, in this figure, grain 19
size
differences between the austenite and the martensite suggest microstructure heterogeneity.

Figure
Figure 4.
4. SEM
SEM images
images for
for all
all the
the analyzed
analyzed zones.
zones. (a)
(a) FZ,
FZ, (b)
(b) fusion
fusion line
line defining
defining the
the border
border between
between
the FZ and the HAZ, and (c) BM.
the FZ and the HAZ, and (c) BM.

The
For aapparition
more detailedof δ-ferrite in the FZcharacterization
microstructure confirms the high temperature
in all the zones, reached
the EBSDduring
maps
the fusion. At first glance, Figure 5a suggests that the δ-ferrite solidified
allowed the phases identification, grain size calculations for all phases, and texture around theevolu-
aus-
tenite grain Figure
tion. Thus, boundaries with an
5a through irregular
Figure skeletal-type
5d present the phasemorphology. This phase
maps overlaid over the stretches
Image
to the limits
Quality (IQ)with
for thethedifferent
HAZ until the fusion
zones, startingline, as shown
with the FZ in Figure
until 5b. From
the BM. this border
For example, in
forward, the bcc crystal structure corresponds to the α’-martensite
the FZ, Figure 5a highlights in blue and red colors the δ-ferrite and γ-austenite phases, due to the defor-
mation-induced
respectively. In this phase transformation
figure, it is observed phenomenon of ASSstructure
that the dendritic during the manufacturing
corresponds to the
process (tube was
primary δ-ferrite phase.supplied in a cold extrusion state). According to Shen et al. [33], the
amountTheofapparition
α’-martensite in a 304L-ASS
of δ-ferrite in the FZisconfirms
a functiontheofhigh
composition,
temperaturestrain, and strain
reached duringrate.
the
TheAtway
fusion. firstthe weldFigure
glance, joint 5ahassuggests
solidified agrees
that with the
the δ-ferrite 𝐶𝑟 ⁄𝑁𝑖around
solidified relationship that
the austenite
controls the phaseswith
grain boundaries present in stainless
an irregular steels, as indicated
skeletal-type morphology.in theThis
Schaeffler
phase diagram.
stretches Thus,
to the
the amount
limits with the 𝐶𝑟 and
of HAZ until𝑁𝑖 can beline,
the fusion evaluated
as shownby the following
in Figure equations
5b. From [34]: forward,
this border
the bcc crystal structure corresponds to the α’-martensite due to the deformation-induced
𝐶𝑟 = %𝐶𝑟 + %𝑀𝑜 + 1.5%𝑆𝑖 + 0.5% 𝑁𝑏 + 2%𝑇𝑖 (3)
phase transformation phenomenon of ASS during the manufacturing process (tube was
supplied in a cold extrusion𝑁𝑖 =state).%𝑁𝑖 + 0.5%𝑀𝑛 to
According + 30%𝐶
Shen et+al.30[33],
%𝑁the 0.06
amount of α’-martensite (4)
in a 304L-ASS is a function of composition, strain, and strain rate.
From
The waythe chemical
the weldcomposition in Tableagrees
joint has solidified 1 and using
with Equations
the Creq /Ni (3) and (4), a value of
eq relationship that
𝐶𝑟 ⁄𝑁𝑖 = 1.67 was obtained. According to Lu et al. [35] and Yazdian et al. [32], if the
controls the phases present in stainless steels, as indicated in the Schaeffler diagram. Thus,
the amount of𝐶𝑟Cr⁄eq𝑁𝑖
relationship and Ni1.25, the FZ will solidify with the presence of ferrite around the
eq can be evaluated by the following equations [34]:
austenite grain boundaries. Given the amounts of 𝐶𝑟 = 20.8 and 𝑁𝑖 = 12.5 for the
studied steel and using Creqthe = Schaeffler
%Cr + %Mo diagram
+ 1.5%Sipresented
+ 0.5% in Nbthe research work of Quitzke
+ 2%Ti (3)
et al. [22], the amount of δ-ferrite should not exceed 10%. This value agrees with the 9%
of δ-ferrite calculated Niby = %Nias
eq EBSD, + indicated
0.5%Mn +in30%C Figure 30(%N − 0.06)
+ 5a. (4)
Other factors also help to keep a δ-ferrite fraction in the FZ. For example, since the δ
→ γ transformation is a diffusion-controlled solid-state transformation, the fast-cooling
rates involved in the GTAW process can be responsible for a non-complete austenite trans-
formation. Nevertheless, δ-ferrite can play a positive role in the ASS’ weld seams because
it acts as a barrier for crack propagation.
Metals 2023, 13,
Metals 2023, 13, 245
245 98 of
of 18
19

Figure 5.
Figure 5. Identification
Identification phase
phase maps
maps overlapped
overlapped on
on the
the image
image quality
quality map
map for
for the
the (a)
(a) FZ,
FZ, (b)
(b) fusion
fusion
line between the FZ and HAZ, (c) BM1, (d) BM2, and (e) grain size distributions for all the phases
line between the FZ and HAZ, (c) BM1, (d) BM2, and (e) grain size distributions for all the phases in
in the analyzed zones.
the analyzed zones.

In zones
From the far from the
chemical fusion lineinand
composition the1HAZ,
Table and usingin theEquations
zone called (3) BM1,
and (4),which corre-
a value of
sponds to the based material that still could be affected by the welding
Creq /Nieq = 1.67 was obtained. According to Lu et al. [35] and Yazdian et al. [32], if the heat, new micro-
structural characteristics
relationship Creq /Nieq > are developed,
1.25, the FZ will as solidify
indicated within Figure 5c. At this
the presence point,around
of ferrite the micro-the
structure presents two phases: the austenite (red color) and the bcc
austenite grain boundaries. Given the amounts of Creq = 20.8 and Nieq = 12.5 for the studied crystal structure cor-
responding
steel and using to the
the Schaeffler
deformation-induced
diagram presented α’-martensite (greenwork
in the research color). Beyond et
of Quitzke theal.BM1
[22],
zone,
the i.e., the
amount ofBM2 zone,should
δ-ferrite Figurenot5d exceed
shows 10%.a mixture
This of austenite
value agreesand with α’-martensite again,
the 9% of δ-ferrite
being the latter
calculated fraction
by EBSD, more dominant.
as indicated in Figure 5a.
The development
Other factors also help of α’-martensite in stainless
to keep a δ-ferrite fraction steels
in the of FZ.
the For
304L series depends
example, since theon δ
→ γ transformation is a diffusion-controlled solid-state transformation, the fast-cooling
microstructural parameters such as the stacking fault energy (𝛾 ). A quick measurement
of thisinvolved
rates parameter can GTAW
in the be conducted
processfromcan the steel chemical
be responsible for composition
a non-complete following
austenitethe
equation proposed
transformation. by Schrammδ-ferrite
Nevertheless, et al. [36]:
can play a positive role in the ASS’ weld seams
because it acts as a barrier for crack propagation.
𝛾 = 53 + 6.2%𝑁𝑖 + 0.7%𝐶𝑟 + 3.2%𝑀𝑛 + 9.3%𝑀𝑜 (5)
In zones far from the fusion line and the HAZ, in the zone called BM1, which cor-
responds
UsingtoEquation
the based(5) material
and thethatBMstill could be
chemical affected byshown
composition the welding
in Table heat, the 𝛾mi-
1, new
crostructural characteristics are developed, as indicated in Figure
value for the studied steel is ~ 15.1 mJ/m , which can be considered low energy according
2 5c. At this point, the
microstructure presents
to Allain et al. [37]. two phases:
Furthermore, theytheestablished
austenite (red for acolor)
highandMn the steelbcc crystal
that structure
twinning was
corresponding to the deformation-induced
the favorite deformation mechanism when the 𝛾
α’-martensiteranges (green
betweencolor).18 Beyond
mJ/m2–45 the BM1
mJ/m 2,

zone,
while i.e., the BM2 zone,
deformation Figure
assisted 5d shows a(dislocation
by dislocation mixture of austenite and α’-martensite
glide) required energies above again,45
being
mJ/m2the latter fraction
. Therefore, more dominant.
the preferred deformation mechanism at room temperature for the
Thesteel
studied development of α’-martensite inphase
is the deformation-induced stainless steels of thei.e.,
transformation, 304L series depends
austenite transforms on
microstructural
into martensite. parameters such as the stacking fault energy (γ SFE ). A quick measurement
of this parameter
Moving on tocan be conducted
grain size evolution frominthethesteel chemical
different zones,composition following the
Figure 5e represents the
equation
grain sizeproposed
distributionsby Schramm et al. indicated
for each zone [36]: in Figure 5a to Figure 4d. As a first obser-
vation, the grain size distributions for the two phases in the FZ show more than one mag-
γSFE = −53 + 6.2%Ni + 0.7%Cr + 3.2%Mn + 9.3%Mo (5)
nitude order difference between the austenite and δ-ferrite average grain size. These dif-
ferences
Using suggest
Equation a heterogeneous
(5) and the BMmicrostructure
chemical compositioninside the shownFZ.inAround
Table 1,thethefusion line,
γSFE value
grain size differences decrease considerably,
2 but the average
for the studied steel is ~15.1 mJ/m , which can be considered low energy according toaustenite grain size is almost
six times
Allain greater
et al. [37]. than the mix of
Furthermore, ferrite
they and martensite.
established for a high These fewerthat
Mn steel differences
twinninginwas the
grain
the size calculations
favorite deformationare related to when
mechanism the consideration
the γSFE ranges of α’-martensite,
between 18 mJ/m which 2 –45possesses
mJ/m2 ,
larger grains than δ-ferrite.
Metals 2023, 13, 245 9 of 18

while deformation assisted by dislocation (dislocation glide) required energies above

45 mJ/m2 . Therefore, the preferred deformation mechanism at room temperature for the
studied steel is the deformation-induced phase transformation, i.e., austenite transforms
into martensite.
Moving on to grain size evolution in the different zones, Figure 5e represents the grain
size distributions for each zone indicated in Figure 5a to Figure 4d. As a first observation,
the grain size distributions for the two phases in the FZ show more than one magnitude
order difference between the austenite and δ-ferrite average grain size. These differences
suggest a heterogeneous microstructure inside the FZ. Around the fusion line, grain size
differences decrease considerably, but the average austenite grain size is almost six times
greater than the mix of ferrite and martensite. These fewer differences in the grain size
calculations are related to the consideration of α’-martensite, which possesses larger grains
than δ-ferrite.
In more distant areas, e.g., BM1 and BM2, the austenitic grain size tends to stabilize
around 8 µm, as illustrated in the OM analysis presented in Figure 3d. In the same zones,
the martensite grain size decreases from 7 µm to 4 µm, and its fraction increases from
22% to 64%. These grain size and phase volume fraction variations can produce overall
grain refinement and material strengthening. Muñoz et al. [38] found a yield strength
improvement in an ASS due to the grain refinement after plastic deformation by increasing
the martensite fraction. Thus, it is clear that the welded tube microstructure owns different
grades of heterogeneity.
The microstructural heterogeneity generated by different phase transformations due
to the GTAW welding and the tube forming process can affect the grain size, the grain
boundary nature, and the dislocation density, as some researchers have proved [39,40]. In
this context, Figure 6 presents the IQ maps together with the different grain boundaries
(black and red colors correspond with the HAGB and LAGB, respectively) and the corre-
sponding misorientation angle distributions. As with the grain sizes, Figure 6a through
Figure 6d also suggest different amounts of LAGB and HAGB depending on the analyzed
zone. The mentioned differences can be interpreted and compared by the misorientation
angle distributions. These distributions demonstrate that the HAGB fractions for all phases
increase as the distance from the FZ also increases, as illustrated in Figure 6e. This analysis
also highlights two things: (1) both phases found the largest LAGB fractions in the FZ, and
(2) in each zone, the phases with bcc crystal structure showed the largest LAGB fractions.
These observations allow seeing microstructure heterogeneity along the zones and
between the phases of each zone. Besides, the most significant fractions of LAGB in
the FZ and its neighborhoods are related to a high density of Geometrically Necessary
Dislocations (GNDs), which controls the curvature of the grain [41–43]. This is because
the more intense thermal gradients produce higher residual stresses around the FZ than
in the BM. On the other hand, metallic materials with two phases, such as duplex, dual
phase, and austenitic stainless steels, present different dislocation densities due to their
different capacities to distribute deformation. According to several authors [38,44,45],
deformation distribution must also be heterogenous between the two phases leading to
additional material hardening. It also demonstrates the creation of plastic gradients around
the weld seam and between the phases of each zone.
 intense thermal gradients produce higher residual stresses around the FZ than in the BM.
On the other hand, metallic materials with two phases, such as duplex, dual phase, and
austenitic stainless steels, present different dislocation densities due to their different ca-
pacities to distribute deformation. According to several authors [38,44,45], deformation
distribution must also be heterogenous between the two phases leading to additional ma-
Metals 2023, 13, 245 10 of 18
terial hardening. It also demonstrates the creation of plastic gradients around the weld
seam and between the phases of each zone.

Figure 6.
Figure 6. IQ
IQ imagens
imagenstogether
togetherwith
withthe
thegrain
grainboundaries
boundariesforfor
thethe
different analyzed
different zones.
analyzed (a) FZ,
zones. (b)
(a) FZ,
fusion line between the FZ and HAZ, (c) BM1, (d) BM2, and (e) misorientation angle distributions
(b) fusion line between the FZ and HAZ, (c) BM1, (d) BM2, and (e) misorientation angle distributions
for all the phases in all the analyzed zones.
for all the phases in all the analyzed zones.

3.2. Texture Heterogeneity

The preferential crystallographic orientation is closely related to microstructural pa-
rameters, such as grain size, dislocation density, and grain boundary orientation, which are
influenced by the solidification, thermal gradients, and the tube-forming process. At first
glance, the highest texture intensities were observed in the FZ, followed by the zones close
to the fusion line and the HAZ, as shown in Figure 7a,b, respectively. However, according
to the (111) pole figure, austenite texture for these zones does not follow any specific
texture component associated with recrystallization, deformation, or shear mechanism.
This is coherent with the solidification process that produced huge austenite grains with
random orientations. Conversely, the (110) pole figure for the δ-ferrite indicates more
symmetric components in both zones. This is because, in the FZ, the δ-ferrite solidified in
the austenite grain boundaries. Furthermore, in the HAZ, texture calculation also consid-
ers the α’-martensite, which conserves its preferential orientation towards the Extrusion
Direction (ED). In zones BM1 and BM2, the pole figures for both phases indicate symmetric
behaviors associated with a monotonic process such as rolling or extrusion, as indicated
in Figure 7c,d. Therefore, the tube’s thermal (solidification and recrystallization) and de-
formation (phase transformation and continuous dynamic recrystallization) history have
favored the apparition of different kinds of textures across the FZ, HAZ, and BM.

3.3. Hardness Evolution

As a consequence of the microstructural changes on the weld seam and its surround-
ings, to determine the joint weak points, the mechanical behavior was analyzed following
hardness profiles in three zones, that is FZ, HAZ, and BM. Thus, dashed lines in Figure 8a
indicate where the hardness profiles were measured. As a first observation, Figure 8b
highlights that the hardness profiles inside the FZ vary significantly over the measured
distance with an average value of 183 HV and a standard deviation value of 10 6 HV. This
behavior certainly is attributed to the heterogeneous dendritic morphology across the FZ, as
indicated in Figure 3b and corroborated by the grain size and grain boundaries evolutions
in Figures 5 and 6, respectively. Then, the HAZ hardness profiles show less scatter and
higher magnitudes than the FZ (average value of 195 HV and a standard deviation value of
6.5 HV) but lower than the BM (average value of 197 HV and a standard deviation value of
5 HV). The hardness differences between the HAZ and the BM are related to the grain size
 close to the fusion line and the HAZ, as shown in Figure 7a and Figure 7b, respectively.
However, according to the 111 pole figure, austenite texture for these zones does not
follow any specific texture component associated with recrystallization, deformation, or
shear mechanism. This is coherent with the solidification process that produced huge aus-
tenite grains with random orientations. Conversely, the 110 pole figure for the δ-ferrite
Metals 2023, 13, 245 indicates more symmetric components in both zones. This is because, in the FZ, the δ-11 of 18
ferrite solidified in the austenite grain boundaries. Furthermore, in the HAZ, texture cal-
culation also considers the α’-martensite, which conserves its preferential orientation to-
andwards the Extrusion
the phase Directionvariations
volume fraction (ED). In zones BM1 and
originating BM2,
from the pole
thermal figures for
gradients. Forboth
example,
phases indicate symmetric behaviors associated with a monotonic process such as rolling
in the BM, the α’-martensite grain size is the smallest, and this phase represents more than
or extrusion, as indicated in Figure 7c and Figure 7d. Therefore, the tube’s thermal (solid-
60%, i.e., the hardest phase is more representative than the softest one, as indicated in
ification and recrystallization) and deformation (phase transformation and continuous
Figure 5d. Therefore, based on the hardness evolution, the weld seam’s weakest points are
dynamic recrystallization) history have favored the apparition of different kinds of tex-
mainly
tureslocated between
across the FZ, HAZ, theand
FZBM.
and the HAZ.

Metals 2023, 13, 245 12 of 19

Figure
Figure 7. Inversepole
7. Inverse polefigure
figure maps
maps and
andthe
thecorresponding pole
corresponding figures
pole for the
figures forfcc
theand
fccbcc
andcrystal
bcc crystal
softest one,(a)asFZ,
structures. indicated inline
(b) fusion Figure 5d. Therefore,
between the FZ andbased
HAZ, on the hardness
(c) BM1, (d) BM2. evolution, the weld
structures. (a) FZ, (b) fusion line between the FZ and HAZ, (c)
seam’s weakest points are mainly located between the FZ and the HAZ. BM1, (d) BM2.
3.3. Hardness Evolution
As a consequence of the microstructural changes on the weld seam and its surround-
ings, to determine the joint weak points, the mechanical behavior was analyzed following
hardness profiles in three zones, that is FZ, HAZ, and BM. Thus, dashed lines in Figure 8a
indicate where the hardness profiles were measured. As a first observation, Figure 8b
highlights that the hardness profiles inside the FZ vary significantly over the measured
distance with an average value of 183 HV and a standard deviation value of 10 6 HV. This
behavior certainly is attributed to the heterogeneous dendritic morphology across the FZ,
as indicated in Figure 3b and corroborated by the grain size and grain boundaries evolu-
tions in Figure 5 and Figure 6, respectively. Then, the HAZ hardness profiles show less
scatter and higher magnitudes than the FZ (average value of 195 HV and a standard de-
viation value of 6.5 HV) but lower than the BM (average value of 197 HV and a standard
deviation value of 5 HV). The hardness differences between the HAZ and the BM are re-
lated to the grain size and the phase volume fraction variations originating from thermal
gradients. For example, in the BM, the α’-martensite grain size is the smallest, and this
phase represents more than 60%, i.e., the hardest phase is more representative than the

Figure 8. (a) hardness profiles localization and (b) hardness measurements on the different zones
Figure 8. (a) hardness profiles localization and (b) hardness measurements on the different zones
and the average curve for each zone (doted lines in figure (a) do not correspond to the exact path
andand
thelocation
averagewhere
curvethe
forhardness
each zone (doted
profiles lines
were in figure
obtained, (a)are
they dojust
notused
correspond to the exact
as an illustration path and
mode).
location where the hardness profiles were obtained, they are just used as an illustration mode).
3.4. Plastic Behavior and Internal Stresses
Uniaxial tensile tests in each zone will be discussed to better understand the micro-
structural heterogeneity impact on the weld seam mechanical response. Figure 9a presents
the engineering (symbol line) and true (line without symbols) stress–strain tensile curves
for the FZ, HAZ, and BM, following the sketch in Figure 2. These curves reflect clear duc-
tility and strength variations with the same tendency of the hardness profiles. So, the FZ
Metals 2023, 13, 245 12 of 18

Metals 2023, 13, 245 3.4. Plastic Behavior and Internal Stresses 13 of 19

Uniaxial tensile tests in each zone will be discussed to better understand the mi-
crostructural heterogeneity impact on the weld seam mechanical response. Figure 9a
microstructure
presents in the HAZ
the engineering and BM
(symbol line)made up of(line
and true austenite
withoutand α’-martensite
symbols) allowstensile
stress–strain good
strength
curves forand
the ductility
FZ, HAZ,in andgood
BM,ratios. In these
following cases,inthe
the sketch α’-martensite
Figure is thereflect
2. These curves phaseclear
that
retards the
ductility andplastic instability.
strength variationsThiswith
is because thetendency
the same α’-martensite
of theishardness
tougher than the austen-
profiles. So, the
ite; reaches
FZ therefore,
thefor a given
lowest plastic deformation,
mechanical strength andthe austenite
ductility will assume
(maximum more plastic
elongation) valuesde-of
formation
all thanwith
three zones the martensite; i.e., austenite
values of ~1100 MPa and will startrespectively.
~45%, to deform plastically
In the HAZ,ahead of mar-
values of
tensite.
~1300 Oneand
MPa example
58% are of found,
hetero-deformation
while the BMisshows
the research workproperties
the highest of Li et al.with
[47] values
in a 304Lof
ASS MPa
1330 with and
a heterogeneous
63% elongation lamella-type microstructure
to failure. These mechanicalobtained
propertiesby confirm
rolling that
and the
subse-
FZ
quent annealing.
presents the poorestIn performance;
this research, however,
the good the
strength–ductility ratio was
mechanical strength associated
and ductility with
values
different
stand outhardening
over otherstages
304L where
ASS with the accumulation of stacking
a 100% austenitic faults/GNDs
microstructure suchandas initiation
the steel
studied by Kumarmartensitic
of strain-induced et al. [46]. transformation took place during the tensile test.

Figure 9.
Figure 9. (a)
(a)True
Trueand
andengineering
engineering tensile
tensile curves,
curves, (b) (b) strain
strain hardening
hardening rate curves,
rate curves, and
and (c) (c) strain
strain maps
maps and profiles for the different analyzed zones before the plastic instability starts (the strain
and profiles for the different analyzed zones before the plastic instability starts (the strain profiles
profiles were obtained along the vertical direction following the pink color arrow, so distance on X
were obtained along the vertical direction following the pink color arrow, so distance on X axis means
axis means the length of the analyzed zone being 0 in the upper part, i.e., the beginning of the ar-
the length of the analyzed zone being 0 in the upper part, i.e., the beginning of the arrow).
row).
The good mechanical performance of the FZ can be associated with the presence of
Phase transformations and heterogeneous microstructures made up of dissimilar
δ-ferrite, which contributes to strength (smaller grain size than austenite) and also acts as
phases with different grain sizes affect the steel hardening capacity, as indicated in Figure
a barrier for the crack propagation in the big austenitic grains, delaying the final fracture.
9b. In this figure, the FZ’s strain hardening rate curve decays fast at the beginning of de-
Additionally, it has been demonstrated by Kain et al. [17] that δ-ferrite helps to improve the
formation,
stress followed
corrosion by short
cracking plateau
strength. Onregions before
the other hand, thethe
plastic instability appears
heterogeneous (plastic
microstructure
instability, according to the Considère criteria [48], takes place when
in the HAZ and BM made up of austenite and α’-martensite allows good strength and the strain hardening
rate at a constant
ductility in good strain
ratios.rate
In equals the flow
these cases, thestress, i.e.,
α’-martensite =
is 𝜎).
the phase that retards the
plastic instability. This is because the α’-martensite is tougher
The fact that the FZ is the first to localize the plastic instability than the
is austenite;
related totherefore,
the large
for a given plastic deformation, the austenite will assume more
free path for the crack propagation inside the austenitic grain size, although plastic deformation
it couldthan
be
the martensite; i.e., austenite will start to deform plastically ahead of
worse without δ-ferrite. On the side of the HAZ and BM, the strain hardening rate curves martensite. One
example
take longerof hetero-deformation
to localize the plastic is the researchThis
instability. workisofbecause,
Li et al. in
[47] in a zones,
these 304L ASS thewith a
strain
hardening rate increases a little before reaching a plateau behavior, so the homogeneous
plastic zone extends beyond a strain of 0.45. In these cases, the heterogeneous microstruc-
ture made up of soft austenite and hard martensite gives rise to heterogeneous
Metals 2023, 13, 245 13 of 18

heterogeneous lamella-type microstructure obtained by rolling and subsequent annealing.

In this research, the good strength–ductility ratio was associated with different hardening
stages where the accumulation of stacking faults/GNDs and initiation of strain-induced
martensitic transformation took place during the tensile test.
Phase transformations and heterogeneous microstructures made up of dissimilar
phases with different grain sizes affect the steel hardening capacity, as indicated in
Figure 9b. In this figure, the FZ’s strain hardening rate curve decays fast at the begin-
ning of deformation, followed by short plateau regions before the plastic instability appears
(plastic instability, according to the Considère criteria [48], takes place when the strain
hardening rate at a constant strain rate equals the flow stress, i.e., dσ dε ε
. = σ).

The fact that the FZ is the first to localize the plastic instability is related to the large free
path for the crack propagation inside the austenitic grain size, although it could be worse
without δ-ferrite. On the side of the HAZ and BM, the strain hardening rate curves take
longer to localize the plastic instability. This is because, in these zones, the strain hardening
rate increases a little before reaching a plateau behavior, so the homogeneous plastic zone
extends beyond a strain of 0.45. In these cases, the heterogeneous microstructure made up
of soft austenite and hard martensite gives rise to heterogeneous deformation states and
plastic gradients that act in favor of good strength–ductility ratios. In this context, in their
review manuscript, Romero-Resendiz et al. [49] explain that due to the presence of phases
with different flow stresses and stacking fault energies, ASSs are ideal candidates to reach a
good compromise between strength and ductility.
The remarkable microstructural variations in grain sizes, grain boundary nature, phase
transformations, and internal stresses across the weld seam and the BM have also influenced
the strain distribution during the tensile tests, as indicated in Figure 9c. These maps
obtained at a deformation just below the plastic instability allow a better understanding of
the welded tube performance. Thus, contrary to what was observed in the tensile curves,
the FZ reaches the largest strain magnitudes of all zones (i.e., strains larger than 1); however,
the deformation was concentrated in a small area, as manifested by the deformation map,
and the peak in the deformation profile. Therefore, the strong strain localization explains
the lack of hardening in the FZ. On the other hand, deformation is evenly distributed in
the HAZ and BM, as shown in the strain maps and profiles. In these zones, the strain
is accumulated in multiple zones before the plastic instability, creating smaller strain
concentrations more evenly spread in all the tensile sample dimensions than the FZ sample.
This behavior can be related to where austenite is located (soft domain) since this phase
experience plastic deformation faster than martensite (hard domain), which possesses a
larger yield strength than austenite.
High internal stresses from thermal gradients and local deformations can also influence
the mechanical response of the elements joined by the welding process. EBSD offers
multiple tools to evaluate the internal stress state based on local misorientations variations.
For instance, Kernel Average Misorientation (KAM) is the average misorientation between
a point on the measurement grid and its neighbors [50]; thus, places with large KAM values
indicate microstructural inhomogeneities.
Figure 10 shows KAM maps and profiles for all the discussed zones. Inside the FZ,
Figure 10a shows continuous KAM variations forming interconnected walls between δ-
ferrite grains; however, KAM values inside δ-ferrite grains are the smallest, producing a
KAM profile with several pointed peaks, as indicated in Figure 10e (the red dashed arrows
indicate the KAM profile orientation). Figure 10b allows analyzing the border between the
FZ and the HAZ, so the KAM map suggests small KAM values on the FZ side and larger
values on the HAZ side. This is due to the lower amount of δ-ferrite observed close to the
fusion line (see Figure 5b), which is attributed to a lower cooling rate than in the middle
of the FZ. Thus, fast solidification kinetics in the middle of the FZ favors keeping δ-ferrite
at room temperature; however, the δ-ferrite fraction is not evenly distributed in the FZ
cross-section, and its fraction decreases towards the fusion line due to a slower solidification
rate. García-García et al. [51] demonstrated the presence of thermal gradients between the
Metals 2023, 13, 245 14 of 18

FZ and the HAZ using the GTAW process on high-Mn steel. They also observed δ-ferrite
Metals 2023, 13, 245 15 of of
nucleation in zones with dendritic morphologies, which agrees with the observations 19
this study.

Figure 10.
Figure 10. KAM
KAM maps.
maps.(a) (a)FZ,
FZ,(b)
(b)fusion
fusionline between
line thethe
between FZFZ
and HAZ,
and (c) (c)
HAZ, BM1, (d)(d)
BM1, BM2, andand
BM2, (e)
KAM profiles for all the analyzed zones (the KAM profiles were obtained following the red dashed
(e) KAM profiles for all the analyzed zones (the KAM profiles were obtained following the red dashed
arrow on each KAM map).
arrow on each KAM map).

Using
On the X-ray diffraction,
HAZ side, the KAM it is possible to
increments arecalculate
producedthe by dislocation density
the α’-martensite and the
formation,
phase volume fraction considering the peak broadening and the
as demonstrated by the KAM profile in Figure 10e. In more distant areas, e.g., the integrated area of BM1
each
diffracted peak. The phases volume fraction and the dislocation densities
zone, the heat transfer impact can still be observed with KAM values similar to the HAZ, obtained with
Equations
as displayed(1)–(2) and the
in Figure diffractograms
10c,e. The behavior in Figure
changes11ain allow for the
the BM2 subsequent
zone where the observa-
KAM
values increase considerably in the α’-martensite, as shown in Figure 10d,e. The and
tions: 1) dislocations densities are higher in the BM than in the FZ; 2) α’-martensite KAM δ-
ferrite are the
increments arephases with with
associated the highest
a highdislocation density corroborating
density of defects, the plastic
such as statistically storedgradi-
and
ents with
GND austenite;which
dislocations, and 3)areδ-ferrite in the FZ
more affected byhas
theahigh
higher dislocation around
temperatures densitythethan austen-
HAZ. On
ite after the welding process.
the FZ side, the high KAM values can also be associated with the shrinkage phenomena
duringAsthe
mentioned before,
solidification thatthe mechanical
gives rise to highproperties
residual of weld seams can be affected by
stresses.
residual stresses originating during the liquid-to-solid shrinkage
Using X-ray diffraction, it is possible to calculate the dislocation during density
solidification
and theor
from the
phase thermal
volume gradients
fraction with the the
considering HAZ.peak Figure 11b indicates
broadening and thetheintegrated
residual stresses
area ofineach
the
FZ and thepeak.
diffracted BM obtained
The phases by X-ray
volume diffraction
fractionafter
and the small sample
dislocation was cut obtained
densities from the tube.
with
So, the sample
Equations is free
(1) and (2) from
and themacroscopic stresses
diffractograms imposed
in Figure 11aduring
allow tube
for thefabrication.
subsequent Theobser-
blue
dots show
vations: (1)the measurement
dislocations placesare
densities either on the
higher external
in the BM thansurfaces
in theorFZ;internally, after cut-
(2) α’-martensite
ting.δ-ferrite
and The values calculated
are the suggest
phases with the that the dislocation
highest FZ withstands compression-type
density corroborating the stresses in
plastic
the middlewith
gradients zone and tension-type
austenite; on the superior
and (3) δ-ferrite in the FZexterior surface.
has a higher This is coherent
dislocation with
density than
the weld seam
austenite shape,
after the where
welding the upper part is wider than the inferior because the melted
process.
material can flow easier in the upper part than in the inferior one, where the melted ma-
terial is more constrained by the tube wall thickness. The same happens in the HAZ,
where the tube resistance to shrinkage during the solidification of the melt pool in the FZ
produces compression and tension-type stresses on the cross and superficial areas, respec-
tively. This behavior has been reported by several researchers that also found significant
tension-type residual stresses in the weld seam surface, then changed to compression type
towards the FZ interior [52,53].
No doubt, welding is a complex manufacturing process where several phase trans-
formations, microstructural changes, and stress states can be involved. Although the stud-
ied ASS suffered multiple changes across the weld seam and BM, its overall strength and
Metals 2023, 13, 245 16 of 19

Metals 2023, 13, 245 15 of 18

ductility remain quite high from a mechanical point of view. However, other properties,
such as corrosion resistance, should be evaluated to fully characterize its performance.

Figure 11. (a) X-ray diffraction patterns from the FZ and the BM zones and (b) residual stresses in
the
the weld
weld seam (values are
seam (values are in
in MPa).
MPa).

4. Conclusions
As mentioned before, the mechanical properties of weld seams can be affected by
residual
The stresses
study oforiginating during
the mechanical themicrostructural
and liquid-to-solid heterogeneity
shrinkage during on asolidification
weld seam ap- or
from the thermal gradients with the HAZ. Figure 11b indicates the residual
plied in an ASS-304L tube through OM, EBSD, X-ray diffraction, hardness measurements, stresses in the
FZ and
and the BM
tensile testsobtained
allowed by X-ray diffraction
drawing the followingafter the small sample was cut from the tube.
conclusions:
So, the sample is free from macroscopic
The GTAW welding process on an ASS resulted stresses imposed during tube
in a biphasic fabrication.inside
microstructure The
blue dots show the measurement places either on the external surfaces
the FZ made up of δ-ferrite (~10 %) and γ-austenite with different dendritic morphologies or internally, after
cutting.
due to aThe values 𝐶𝑟
favorable calculated
⁄𝑁𝑖 =suggest that the FZδ-ferrite
1.67. Moreover, withstands compression-type
created interconnected stresses in
disloca-
the middle
tion walls inzone and as
the FZ, tension-type
illustrated byon KAM
the superior
analysis.exterior surface.
This gave rise This is coherent
to strength with
improve-
the weld seam shape, where the upper part is wider than the inferior because
ments because of the δ-ferrite grain size (5.9 µm) smaller than austenite (253 µm) and the the melted
material can flow easier in the upper part than in the inferior one, where the melted material
obstacles (interconnected dislocation walls) for the crack’s free propagation in the austen-
is more constrained by the tube wall thickness. The same happens in the HAZ, where the
ite large grains.
tube resistance to shrinkage during the solidification of the melt pool in the FZ produces
Thermal gradients produced grain growth from the fusion line to the BM with values
compression and tension-type stresses on the cross and superficial areas, respectively. This
ranging from 100 µm to 10 µm. It was also found higher dislocation density and smaller
behavior has been reported by several researchers that also found significant tension-type
grain size in the α’-martensite (𝜌 = 4.11 × 1014 m−2) than in austenite (𝜌 = 1.92 × 1014 m−2).
residual stresses in the weld seam surface, then changed to compression type towards the
However, the mechanical properties between the HAZ and the BM only showed 30 MPa
FZ interior [52,53].
strength and 5% elongation differences.
No doubt, welding is a complex manufacturing process where several phase transfor-
Deformation maps demonstrated that the lowest strength and ductility in the FZ (ul-
mations, microstructural changes, and stress states can be involved. Although the studied
timate
ASS tensile strength
suffered multipleofchanges
1100 MPa and maximum
across the weld seam elongation
and BM, of 45%) was associated
its overall strengthwith
and
aductility
strong localized
remain quite high from a mechanical point of view. However, otherinproperties,
strain allowing the plastic instability occurs earlier than the HAZ
(ultimate tensile strength
such as corrosion of 1300
resistance, MPabeand
should maximum
evaluated elongation
to fully of 58%)
characterize itsand BM (ultimate
performance.
tensile strength of 1330 MPa and maximum elongation of 63%), where the strain localized
in multiple points delaying the plastic instability.
4. Conclusions
The study of the mechanical and microstructural heterogeneity on a weld seam applied
in an ASS-304L tube through OM, EBSD, X-ray diffraction, hardness measurements, and
tensile tests allowed drawing the following conclusions:
Metals 2023, 13, 245 16 of 18

The GTAW welding process on an ASS resulted in a biphasic microstructure inside the
FZ made up of δ-ferrite (~10%) and γ-austenite with different dendritic morphologies due to
a favorable Creq /Nieq = 1.67. Moreover, δ-ferrite created interconnected dislocation walls in
the FZ, as illustrated by KAM analysis. This gave rise to strength improvements because of the
δ-ferrite grain size (5.9 µm) smaller than austenite (253 µm) and the obstacles (interconnected
dislocation walls) for the crack’s free propagation in the austenite large grains.
Thermal gradients produced grain growth from the fusion line to the BM with
values ranging from 100 µm to 10 µm. It was also found higher dislocation density
and smaller grain size in the α’-martensite (ρα0 = 4.11 × 1014 m−2 ) than in austenite
(ργ = 1.92 × 1014 m−2 ). However, the mechanical properties between the HAZ and the BM
only showed 30 MPa strength and 5% elongation differences.
Deformation maps demonstrated that the lowest strength and ductility in the FZ
(ultimate tensile strength of 1100 MPa and maximum elongation of 45%) was associated
with a strong localized strain allowing the plastic instability occurs earlier than in the HAZ
(ultimate tensile strength of 1300 MPa and maximum elongation of 58%) and BM (ultimate
tensile strength of 1330 MPa and maximum elongation of 63%), where the strain localized
in multiple points delaying the plastic instability.
After all the microstructural, texture, and residual stress variations across the weld
seam and the BM, from a mechanical point of view, the FZ is revealed as the weakest
point of the weld seam. However, this zone reached a mechanical strength higher than 1
GPa with a homogeneous deformation zone of 0.4 true strain, which can be considered a
good strength–ductility ratio. This performance was highly related to the heterogeneous
structure in the FZ originated during the solidification process.

Author Contributions: Conceptualization, J.A.M.; methodology, J.A.M., E.D., M.A. and R.B.; soft-
ware, J.A.M.; validation, R.B. and J.M.C.; formal analysis, J.A.M. and E.D.; investigation, J.A.M., V.T.
and P.R.; resources, M.A., R.B. and J.M.C.; data curation, J.A.M., V.T. and P.R.; writing—original draft
preparation, J.A.M. and J.M.C.; writing—review and editing, J.A.M., R.B. and J.M.C.; supervision,
R.B. and J.M.C.; project administration, R.B. and J.M.C.; funding acquisition, M.A., R.B. and J.M.C.
All authors have read and agreed to the published version of the manuscript.
Funding: This research received no external funding.
Data Availability Statement: Not applicable.
Acknowledgments: The authors thank Alexander Komissarov from National University of Science
and Technology “MISIS” for providing the studied tubes.
Conflicts of Interest: The authors declare no conflict of interest.

References
1. Mohan Kumar, S.; Siva Shanmugam, N. Effect of Heat Input and Weld Chemistry on Mechanical and Microstructural Aspects of
Double Side Welded Austenitic Stainless Steel 321 Grade Using Tungsten Inert Gas Arc Welding Process. Materwiss. Werksttech.
2020, 51, 349–367. [CrossRef]
2. Costanza, G.; Sili, A.; Tata, M.E. Weldability of Austenitic Stainless Steel by Metal Arc Welding with Different Shielding Gas.
Procedia Struct. Integr. 2016, 2, 3508–3514. [CrossRef]
3. Sirohi, S.; Pandey, C.; Goyal, A. Role of Heat-Treatment and Filler on Structure-Property Relationship of Dissimilar Welded Joint
of P22 and F69 Steel. Fusion Eng. Des. 2020, 159, 111935. [CrossRef]
4. Rathod, D.W. Chapter 4—Comprehensive Analysis of Gas Tungsten Arc Welding Technique for Ni-Base Weld Overlay. In Advanced
Welding and Deforming; Paulo Davim, J., Gupta, K., Gupta, K., Paulo Davim, J., Eds.; Handbooks in Advanced Manufacturing;
Elsevier: Amsterdam, The Netherlands, 2021; pp. 105–126. ISBN 978-0-12-822049-8.
5. Smith, P. Chapter 6—Fabrication, Assembly, and Erection. In The Fundamentals of Piping Design; Smith, P., Ed.; Gulf Publishing
Company: Houston, TX, USA, 2007; pp. 171–189. ISBN 978-1-933762-04-3.
6. Komissarov, A.A.; Sokolov, P.Y.; Tikhonov, S.M.; Sidorova, E.P.; Mishnev, P.A.; Matrosov, M.Y.; Kuznetsov, D. V Production of
Low-Carbon Steel Sheet for Oil-Industry Pipe. Steel Transl. 2018, 48, 748–753. [CrossRef]
7. Chuaiphan, W.; Srijaroenpramong, L. Effect of Hydrogen in Argon Shielding Gas for Welding Stainless Steel Grade SUS 201 by
GTA Welding Process. J. Adv. Join. Process. 2020, 1, 100016. [CrossRef]
8. Zhu, Z.; Ma, X.; Wang, C.; Mi, G.; Zheng, S. The Metallurgical Behaviors and Crystallographic Characteristic on Macro
Deformation Mechanism of 316 L Laser-MIG Hybrid Welded Joint. Mater. Des. 2020, 194, 108893. [CrossRef]
Metals 2023, 13, 245 17 of 18

9. Ogawa, T.; Koseki, T. Weldability of Newly Developed Austenitic Alloys for Cryogenic Service: Part 1—Up-to-Date Overview of
Welding Technology. Weld. J. 1987, 66, 8–17.
10. Padilha, A.F.; Rios, P.R. Decomposition of Austenite in Austenitic Stainless Steels. ISIJ Int. 2002, 42, 325–327. [CrossRef]
11. Zhang, S.; Wang, P.; Li, D.; Li, Y. Investigation of the Evolution of Retained Austenite in Fe–13%Cr–4%Ni Martensitic Stainless
Steel during Intercritical Tempering. Mater. Des. 2015, 84, 385–394. [CrossRef]
12. Yin, Y.; Faulkner, R.; Starr, F. 5—Austenitic Steels and Alloys for Power Plants. In Structural Alloys for Power Plants; Shirzadi,
A., Jackson, S., Eds.; Woodhead Publishing Series in Energy; Woodhead Publishing: Sawston, UK, 2014; pp. 105–152. ISBN
978-0-85709-238-0.
13. Was, G.S.; Ukai, S. Chapter 8—Austenitic Stainless Steels. In Structural Alloys for Nuclear Energy Applications; Odette, G.R., Zinkle,
S.J., Eds.; Elsevier: Boston, MA, USA, 2019; pp. 293–347. ISBN 978-0-12-397046-6.
14. Michler, T. Austenitic Stainless Steels. In Reference Module in Materials Science and Materials Engineering; Elsevier: Amsterdam, The
Netherlands, 2016; ISBN 978-0-12-803581-8.
15. Tedmon, C.S.; Vermilyea, D.A.; Rosolowski, J.H. Intergranular Corrosion of Austenitic Stainless Steel. J. Electrochem. Soc. 1971,
118, 192. [CrossRef]
16. Fregonese, M.; Idrissi, H.; Mazille, H.; Renaud, L.; Cetre, Y. Initiation and Propagation Steps in Pitting Corrosion of Austenitic
Stainless Steels: Monitoring by Acoustic Emission. Corros. Sci. 2001, 43, 627–641. [CrossRef]
17. Kain, V. 5—Stress Corrosion Cracking (SCC) in Stainless Steels. In Stress Corrosion Cracking; Raja, V.S., Shoji, T., Eds.; Woodhead
Publishing Series in Metals and Surface Engineering; Woodhead Publishing: Sawston, UK, 2011; pp. 199–244. ISBN 978-1-84569-
673-3.
18. Böhner, A.; Niendorf, T.; Amberger, D.; Höppel, H.W.; Göken, M.; Maier, H.J. Martensitic Transformation in Ultrafine-Grained
Stainless Steel AISI 304L Under Monotonic and Cyclic Loading. Metals 2012, 2, 56–64. [CrossRef]
19. Hamada, A.S.; Karjalainen, L.P.; Surya, P.K.C.V.; Misra, R.D.K. Fatigue Behavior of Ultrafine-Grained and Coarse-Grained Cr–Ni
Austenitic Stainless Steels. Mater. Sci. Eng. A 2011, 528, 3890–3896. [CrossRef]
20. Lima, A.S.; Nascimento, A.M.; Abreu, H.F.G.; de Lima-Neto, P. Sensitization Evaluation of the Austenitic Stainless Steel AISI
304L, 316L, 321 and 347. J. Mater. Sci. 2005, 40, 139–144. [CrossRef]
21. Muthupandi, V.; Bala Srinivasan, P.; Seshadri, S.K.; Sundaresan, S. Effect of Weld Metal Chemistry and Heat Input on the Structure
and Properties of Duplex Stainless Steel Welds. Mater. Sci. Eng. A 2003, 358, 9–16. [CrossRef]
22. Quitzke, C.; Schröder, C.; Mandel, M.; Krüger, L.; Volkova, O.; Wendler, M. Solidification of Plasma TIG-Welded N-Alloyed
Austenitic CrMnNi Stainless Steel. Weld. World 2022, 66, 2217–2229. [CrossRef]
23. Shankar, V.; Gill, T.P.S.; Mannan, S.L.; Sundaresan, S. Solidification Cracking in Austenitic Stainless Steel Welds. Sadhana 2003, 28,
359–382. [CrossRef]
24. Igathinathane, C.; Pordesimo, L.O.; Columbus, E.P.; Batchelor, W.D.; Methuku, S.R. Shape Identification and Particles Size
Distribution from Basic Shape Parameters Using ImageJ. Comput. Electron. Agric. 2008, 63, 168–182. [CrossRef]
25. Bachmann, F.; Hielscher, R.; Schaeben, H. Texture Analysis with MTEX–Free and Open Source Software Toolbox. Solid State
Phenom. 2010, 160, 63–68. [CrossRef]
26. Muñoz, J.A.; Komissarov, A. Back Stress and Strength Contributions Evolution of a Heterogeneous Austenitic Stainless Steel
Obtained after One Pass by Equal Channel Angular Sheet Extrusion (ECASE). Int. J. Adv. Manuf. Technol. 2020, 109, 607–617.
[CrossRef]
27. De, A.K.; Murdock, D.C.; Mataya, M.C.; Speer, J.G.; Matlock, D.K. Quantitative Measurement of Deformation-Induced Martensite
in 304 Stainless Steel by X-Ray Diffraction. Scr. Mater. 2004, 50, 1445–1449. [CrossRef]
28. Lutterotti, L.; Matthies, S.; Wenk, H.R. MAUD (Material Analysis Using Diffraction): A User Friendly Java Program for Rietveld
Texture Analysis and More. In Proceedings of the Twelfth International Conference on Textures of Materials (ICOTOM-12),
Montreal, QC, Canada, 9–13 August 1999; NRC Research Press: Ottawa, ON, Canada, 1999; p. 1599.
29. Zhao, Y.H.; Sheng, H.W.; Lu, K. Microstructure Evolution and Thermal Properties in Nanocrystalline Fe during Mechanical
Attrition. Acta Mater. 2001, 49, 365–375. [CrossRef]
30. Ghani, A.F.A.; Ali, M.B.; Dharmalingam, S.; Mahmud, J. Digital Image Correlation (DIC) Technique in Measuring Strain Using
Opensource Platform Ncorr. J. Adv. Res. Appl. Mech. 2016, 26, 10–21. [CrossRef]
31. Ostovan, F.; Shafiei, E.; Toozandehjani, M.; Mohamed, I.F.; Soltani, M. On the Role of Molybdenum on the Microstructural,
Mechanical and Corrosion Properties of the GTAW AISI 316 Stainless Steel Welds. J. Mater. Res. Technol. 2021, 13, 2115–2125.
[CrossRef]
32. Yazdian, N.; Mohammadpour, M.; Razavi, R.; Kovacevic, R. Hybrid Laser/Arc Welding of 304L Stainless Steel Tubes, Part 2–Effect
of Filler Wires on Microstructure and Corrosion Behavior. Int. J. Press. Vessel. Pip. 2018, 163, 45–54. [CrossRef]
33. Shen, Y.F.; Li, X.X.; Sun, X.; Wang, Y.D.; Zuo, L. Twinning and Martensite in a 304 Austenitic Stainless Steel. Mater. Sci. Eng. A
2012, 552, 514–522. [CrossRef]
34. Suutala, N.; Takalo, T.; Moisio, T. The Relationship between Solidification and Microstructure in Austenitic and Austenitic-Ferritic
Stainless Steel Welds. Metall. Trans. A 1979, 10, 512–514. [CrossRef]
35. Lu, B.T.; Chen, Z.K.; Luo, J.L.; Patchett, B.M.; Xu, Z.H. Pitting and Stress Corrosion Cracking Behavior in Welded Austenitic
Stainless Steel. Electrochim. Acta 2005, 50, 1391–1403. [CrossRef]
Metals 2023, 13, 245 18 of 18

36. Schramm, R.E.; Reed, R.P. Stacking Fault Energies of Seven Commercial Austenitic Stainless Steels. Metall. Trans. A 1975, 6, 1345.
[CrossRef]
37. Allain, S.; Chateau, J.-P.; Bouaziz, O. A Physical Model of the Twinning-Induced Plasticity Effect in a High Manganese Austenitic
Steel. Mater. Sci. Eng. A 2004, 387–389, 143–147. [CrossRef]
38. Muñoz, J.A.; Komissarov, A.; Avalos, M.; Bolmaro, R.E. Mechanical and Microstructural Behavior of a Heterogeneous Austenitic
Stainless Steel Processed by Equal Channel Angular Sheet Extrusion (ECASE). Mater. Sci. Eng. A 2020, 792, 139779. [CrossRef]
39. Muñoz, J.A.; Komissarov, A.; Mejía, I.; Hernández-Belmontes, H.; Cabrera, J.-M. Characterization of the Gas Tungsten Arc
Welding (GTAW) Joint of Armco Iron Nanostructured by Equal-Channel Angular Pressing (ECAP). J. Mater. Process. Technol. 2021,
288, 116902. [CrossRef]
40. Muñoz, J.A.; Cobos, O.F.H.; M’Doihoma, R.; Avalos, M.; Bolmaro, R.E. Inducing Heterogeneity in an Austenitic Stainless Steel by
Equal Channel Angular Sheet Extrusion (ECASE). J. Mater. Res. Technol. 2019, 8, 2473–2479. [CrossRef]
41. Nye, J.F. Some Geometrical Relations in Dislocated Crystals. Acta Metall. 1953, 1, 153–162. [CrossRef]
42. Muñoz, J.A.; Avalos, M.; Schell, N.; Brokmeier, H.G.; Bolmaro, R.E. Comparison of a Low Carbon Steel Processed by Cold Rolling
(CR) and Asymmetrical Rolling (ASR): Heterogeneity in Strain Path, Texture, Microstructure and Mechanical Properties. J. Manuf.
Process. 2021, 64, 557–575. [CrossRef]
43. Muñoz, J.A. Geometrically Necessary Dislocations (GNDs) in Iron Processed by Equal Channel Angular Pressing (ECAP). Mater.
Lett. 2019, 238, 42–45. [CrossRef]
44. Muñoz, J.A.; Chand, M.; Signorelli, J.W.; Calvo, J.; Cabrera, J.M. Strengthening of Duplex Stainless Steel Processed by Equal
Channel Angular Pressing (ECAP). Int. J. Adv. Manuf. Technol. 2022, 123, 2261–2278. [CrossRef]
45. Ramazani, A.; Mukherjee, K.; Schwedt, A.; Goravanchi, P.; Prahl, U.; Bleck, W. Quantification of the Effect of Transformation-
Induced Geometrically Necessary Dislocations on the Flow-Curve Modelling of Dual-Phase Steels. Int. J. Plast. 2013, 43, 128–152.
[CrossRef]
46. Kumar, S.S.S.; Vasanth, M.; Singh, V.; Ghosal, P.; Raghu, T. An Investigation of Microstructural Evolution in 304L Austenitic
Stainless Steel Warm Deformed by Cyclic Channel Die Compression. J. Alloys Compd. 2017, 699, 1036–1048. [CrossRef]
47. Li, J.; Fang, C.; Liu, Y.; Huang, Z.; Wang, S.; Mao, Q.; Li, Y. Deformation Mechanisms of 304L Stainless Steel with Heterogeneous
Lamella Structure. Mater. Sci. Eng. A 2019, 742, 409–413. [CrossRef]
48. Considère, A. Memoire Sur l’emploi Du Fer et de l’acier Dans Les Constructions. Ann. Ponts Chaussées Sem. 1885, 9, 574–775.
49. Romero-Resendiz, L.; El-Tahawy, M.; Zhang, T.; Rossi, M.C.; Marulanda-Cardona, D.M.; Yang, T.; Amigó-Borrás, V.; Huang, Y.;
Mirzadeh, H.; Beyerlein, I.J.; et al. Heterostructured Stainless Steel: Properties, Current Trends, and Future Perspectives. Mater.
Sci. Eng. R Rep. 2022, 150, 100691. [CrossRef]
50. Wright, S.I.; Nowell, M.M.; Field, D.P. A Review of Strain Analysis Using Electron Backscatter Diffraction. Microsc. Microanal.
2011, 17, 316–329. [CrossRef] [PubMed]
51. García-García, V.; Mejía, I.; Hernández-Belmontes, H.; Reyes-Calderón, F.; Benito, J.A.; Cabrera, J.M. Computational Simulation of
Thermo-Mechanical Field and Fluid Flow and Their Effect on the Solidification Process in TWIP Steel Welds. J. Manuf. Process.
2022, 84, 100–120. [CrossRef]
52. Ishigami, A.; Roy, M.J.; Walsh, J.N.; Withers, P.J. The Effect of the Weld Fusion Zone Shape on Residual Stress in Submerged Arc
Welding. Int. J. Adv. Manuf. Technol. 2017, 90, 3451–3464. [CrossRef]
53. Ganesh, K.C.; Vasudevan, M.; Balasubramanian, K.R.; Chandrasekhar, N.; Mahadevan, S.; Vasantharaja, P.; Jayakumar, T.
Modeling, Prediction and Validation of Thermal Cycles, Residual Stresses and Distortion in Type 316 LN Stainless Steel Weld
Joint Made by TIG Welding Process. Procedia Eng. 2014, 86, 767–774. [CrossRef]

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