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Deformation mechanisms in meta-stable and nitrogen-stabilized austenitic stainless
steel during severe surface deformation
Wang, Bo; Hong, Chuanshi; Winther, Grethe; Christiansen, Thomas L.; Somers, Marcel A. J.
Published in:
Materialia
Link to article, DOI:
10.1016/j.mtla.2020.100751
Publication date:
2020
Document Version
Publisher's PDF, also known as Version of record
Link back to DTU Orbit
Citation (APA):
Wang, B., Hong, C., Winther, G., Christiansen, T. L., & Somers, M. A. J. (2020). Deformation mechanisms in
meta-stable and nitrogen-stabilized austenitic stainless steel during severe surface deformation. Materialia, 12,
[100751]. https://doi.org/10.1016/j.mtla.2020.100751
Materialia 12 (2020) 100751
Contents lists available at ScienceDirect
Materialia
journal homepage: www.elsevier.com/locate/mtla
Full Length Article
Deformation mechanisms in meta-stable and nitrogen-stabilized austenitic
stainless steel during severe surface deformation
Bo Wang
∗
, Chuanshi Hong , Grethe Winther , Thomas L. Christiansen , Marcel A.J. Somers
Department of Mechanical Engineering, Technical University of Denmark, DK-2800 Kongens Lyngby, Denmark
Keywords:
AISI 304L stainless steel
High-temperature solution nitriding
Surface roller burnishing
Deformation-induced martensitic
transformation
Austenitic nanocrystallites
AISI 304L stainless steel in austenitized and in solution nitrided condition was severely mechanically deformed
by surface roller burnishing. High-temperature solution nitriding was applied to achieve a nitrogen-concentration
depth prole, leading to a depth-gradient in the austenite stability. X-ray diraction, electron microscopy and
hardness indentation were applied for characterization of the graded microstructures obtained by combining a
composition prole and a deformation prole. While severe plastic surface straining of an austenitized speci-
men leads to a deformation-induced transformation of austenite into martensite, the solution nitrided specimen
remains austenitic upon deformation, even in the region where nanocrystallization occurs. The deformation mech-
anisms operable in the nitrogen-stabilized austenitic stainless steel, i.e. twinning or dislocation glide, depend on
the combination of applied plastic strain/strain rate, and the nitrogen-concentration dependent stacking fault
energy.
1.
Introduction
Metastable austenitic stainless steels are materials of high interest
for a plethora of engineering applications due to their intrinsically high
corrosion resistance and good formability [1] . One of the most e-
cient ways to achieve reasonable strength and fatigue resistance in these
materials is severe plastic surface deformation (SPSD). Conventionally,
SPSD, such as ultrasonic shot peening, deep rolling and surface me-
chanical attrition treatment (SMAT), leads to signicant microstructural
changes associated with grain renement to submicron-/nano-meter di-
mensions, formation of micron-scale defects and deformation-induced
martensitic transformation [2–5] . Despite a positive contribution to
strength, the presence of deformation-induced martensite (i) negatively
inuences the resistance against aqueous corrosion; (ii) increases the
susceptibility to hydrogen embrittlement; (iii) results in a reduction in
the ductility and (iv) is particularly harmful during low-temperature sur-
face nitriding [6–11] . Therefore, the prevention of deformation-induced
martensite upon SPSD is necessary for metastable austenitic stainless
steels subjected to forming applications.
Nitrogen is a potent austenite stabilizer and oers benecial eects
for austenitic stainless steels, involving enhanced yield strength without
signicant reduction in ductility, improved electrochemical properties
and a reduction of the nickel content as austenite stabilizer [12–14] .
Nitrogen can be introduced to stainless steels in the liquid state to
manufacture a so-called high nitrogen steel or in the solid state at
∗
Corresponding author.
E-mail address: bwang@mek.dtu.dk (B. Wang).
elevated temperature by thermochemical treatment in a nitrogen-
containing gaseous atmosphere. Both procedures are viable means to
stabilize austenite from transforming into martensite on deformation
and both are applied in industrial practice. Dissolution of nitrogen into
the liquid phase is associated with a relatively low maximum nitrogen
solubility and requires high pressure electroslag remelting [13–15] .
Dissolution in the solid state can be achieved by high-temperature so-
lution nitriding (HTSN), and is limited in the depth range [16 , 17] . The
achievable case depth is several hundred microns (up to a few mm’s)
deep and depends on the process conditions. The nitrogen content at
the surface reects equilibrium between gas and solid and depends
on the alloy composition, the applied temperature and the (total or
partial) N
2
pressure in the gas [11] . In previous research activities,
HTSN was utilized for suppressing the martensitic transformation
during subsequent cold drawing [10] or rolling [18 , 19] of type AISI
304 steel. These investigations conrmed that eective stabilization
of austenite can be achieved by introducing nitrogen in solid solution.
This suggests the possibility of severe deformation of austenite that has
been stabilized by prior HTSN, and thereby potentially the development
of a nano-crystalline surface layer in austenite.
The eect of nitrogen on deformation mechanisms in austenite is
likely to be associated with a modication of the stacking fault en-
ergy (SFE), which plays a key role in the operable deformation mech-
anism in austenitic stainless steels - phase transformation, twinning, or
dislocation glide by slip and the associated strain hardening behavior
https://doi.org/10.1016/j.mtla.2020.100751
Received 15 April 2020; Accepted 28 May 2020
Available online 1 June 2020
2589-1529/© 2020 Acta Materialia Inc. Published by Elsevier B.V. This is an open access article under the CC BY license.
( http://creativecommons.org/licenses/by/4.0/ )
B. Wang, C. Hong and G. Winther et al. Materialia 12 (2020) 100751
[20–22] . The dependence of SFE on the nitrogen content is not accu-
rately known, as conicting values can be found in the scientic liter-
ature [10 , 20 , 21 , 23] . It is anticipated that a variation of nitrogen con-
tent with depth, as obtained after HTSN, is associated with a depth-
distribution of SFE in the solution nitrided case. This, coupled with a
depth-gradient in the degree of deformation and the strain rate as in-
duced by SPSD, provides a means to track the microstructure character-
istics at dierent levels of plastic strain, strain rate, and SFE.
In the present work, a commercial AISI 304L stainless steel is HTSN
treated and subsequently severely deformed by surface roller burnishing
(SRB). The objectives of the present study, are (i) to establish a nitrogen-
concentration depth prole and associated gradient in the austenite sta-
bility prior to SRB, (ii) to interpret the austenite stabilization and mi-
crostructure evolution at dierent levels of strain, strain rate, and SFE,
and (iii) to develop a novel approach to obtain a graded nanostruc-
tured/ultrane surface layer without martensite formation.
2.
Experimental procedure
The starting material AISI 304L stainless steel was provided as plates
with size of 70 mm × 30 mm × 3 mm in annealed condition. The chem-
ical composition of the alloy is (in wt %): 0.023 C, 18.35 Cr, 8.51 Ni,
0.41 Si, 0.83 Mn, 0.03 P, 0.004 S and balance Fe; the SFE for an alloy
with this composition is ~18 mJ/m
2
[20] .
The specimens were solution nitrided at 1150 °C (1423 K) for 2 h in
0.5 bar N
2
(total pressure). As a reference one as-received plate was sub-
jected to a solution heat treatment (SHT), i.e. an austenitization, at the
same temperature and for the same time in argon at atmospheric pres-
sure. Both treatments were terminated by high-pressure gas quenching
to avoid austenite decomposition and/or the formation of nitride precip-
itates during cooling. Subsequently, the surfaces of the SHT and HTSN
specimens were severely plastically deformed through SRB based on a
pin-on-disk conguration similar to that described in Ref. [24] . During
SRB, the roller was pressed against the surface of a steel plate mounted
on a turning lathe, while the steel plate rotated at a speed of 300 rpm.
The roller was made of high speed steel and has a hardness of HV 720–
772.
The roller is 8 mm wide and has a radius of curvature of 6 mm.
The roller rolled onto the surface of the plate and was simultaneously
traveled back and forth at 0.038 mm/revolution in the radial direction
of the rotating plate. Thus, a circular region of 30 mm in diameter was
processed on each specimen. In order to obtain dierent degrees of de-
formation, two feed depths, i.e. the maximum depth of the roller into
the plate, were applied: 300 µm and 500 µm. A surface roughness of
S
a
= 0.3 µm was achieved under the applied SRB feed depth range. Met-
alworking water soluble coolant Quaker cool 7350 BFF was applied for
cooling and lubricating during the SRB process.
Glow discharge optical emission spectroscopy (GD-OES by Horiba
Jobin Yvon GD proler 2) was applied to determine the nitrogen con-
centration prole after HTSN; depth-proling was achieved in a plasma
at 1000 Pa and 40 W. Since the depth range of nitrogen is beyond the
depth range of accurate GD-OES depth proling, the nitrogen concen-
tration prole was measured using spot measurements at multiple lo-
cations, starting from dierent depths on the specimen as realized by
successively removing thin layers by gentle polishing. Quantication of
the nitrogen concentration prole was achieved using reference materi-
als, including a nitrogen-free AISI 304L stainless steel and a JK 49 steel
with a nitrogen content of 2 wt %. Thermodynamic equilibrium calcu-
lations related to solution nitriding of the alloys were performed using
Thermo-Calc 2017b with the TCFE Steels/Fe-Alloys v6.2 database.
A Bruker D8 AXS X-ray diractometer with Cr K 𝛼 radiation was used
to identify the crystalline phases in surface or subsurface of the treated
specimens. To study the structures along the depth direction using XRD
investigation, the specimen was examined at dierent depths by succes-
sively electropolishing o thin layers of 19–30 µm. The cross-sections of
the specimens were mechanically polished and etched at room temper-
ature with Kalling’s reagent for metallographic investigation. Reected
light microscopy was carried out on the cross-sections with a Zeiss Jena
Neophot 30 microscope to either observe the nitrided layer or dieren-
tiate the morphological features of the deformed microstructure treated
at dierent feed depths. The microstructural evolution along sections
perpendicular to the treated surface of the SRB specimens was further
investigated using a combination of electron channeling contrast imag-
ing (ECCI) and electron backscatter diraction (EBSD) on a Zeiss Supra
35 scanning electron microscope (SEM). ECCI was performed at 15 kV
acceleration voltage and a working distance of 6–7 mm, depending on
the specimen shape. For the EBSD procedure, the cross-section of the
specimen was electropolished in a ‘‘Struers Electrolyte A2’’ solution at
25 V for ~30 s in a LectroPol-5 electrolytic polishing machine. EBSD
analysis was performed on a selection of regions with a step size of 30
nm using an accelerating voltage of 20 kV. A thin foil with size of 8
µm × 9 µm was extracted perpendicularly to the feed direction in the
outermost surface region by a focused ion beam (FIB) facility (Tescan
Lyra 3 GMU) to achieve a reliable observation of the ultrane structure.
The specimen thus prepared was welded onto a platinum clamp and ex-
amined in an FEI Tecnai T20 Transmission Electron Microscope (TEM),
operated at an accelerating voltage of 120 kV. The hardness distribution
in the nitrided/deformed surface layers was measured using a Future-
Tech FM-700 hardness tester applying a load of 25 g at a dwelling time
of 10 s at various depths. Presented hardness values are the average of
5 indentations.
3. Results
3.1. Solution nitriding treatment
The eect of solution nitriding was veried by comparison of the HT-
SNed microstructure and the nitrogen concentration prole with those
in as-annealed 304L specimen by means of thermodynamic calcula-
tion and experimental evaluation. An isopleth showing the evolution of
phase stability with nitrogen content, displays the 𝛾-phase eld in the
temperature range 800 °C to 1350 °C in Fig. 1 (a). Iso activity lines as
referred to by the given N
2
pressures are superimposed on the isopleth,
so the equilibrium nitrogen content is given as a function of tempera-
ture and N
2
pressure.
1
HTSN conditions of 1150 °C and 0.5 bar were
selected such that the 𝛾 -phase is stable for the entire composition range
from 0.42 wt % at the surface, in equilibrium with the N
2
gas, to the
nitrogen-free core. No chromium nitrides are expected to develop dur-
ing nitrogen dissolution. The experimental nitrogen concentration pro-
le along the diusion direction is shown in Fig. 1 b. After 2 hours of
HTSN at 1150 °C, the absorbed nitrogen atoms have diused to a depth
of ~ 450 µm and the nitrogen content averaged over the rst 2 µm is
0.40 wt %, in good agreement with the nitrogen content predicted by
assuming thermodynamic equilibrium at the surface ( Fig. 1 a).
The grain size on the right-hand side of the micrograph in Fig. 1 d
reects the eect of temperature on grain size, while some of the grains
in the nitrogen-containing region to the left are appreciably larger.
Evidently, substantial grain growth (from ~21 to ~87 µm) occurred
along with the dissolution of nitrogen in the solid state at 1150 °C (cf.
Fig. 1 c and d). A cross-sectional micrograph of the solution nitrided
specimen ( Fig. 1 d) shows no indications of precipitate formation
at grain boundaries or in the grain interior. As long as the nitrides
are avoided the nitrogen uptake through HTSN is benecial for the
electrochemical properties [25] , which also leads to slower etching
rates close to the surface compared with that in the core; hence, the
faint microstructural details in the surface-adjacent zone.
1
Strictly speaking this is the partial pressure of N
2
at 1 bar total pressure.
However, since the total pressure has a negligible eect on the position of the
phase boundaries in Fig. 1 a for the pressure range considered (up to 1 bar), total
and partial pressure are interchangeable.
B. Wang, C. Hong and G. Winther et al. Materialia 12 (2020) 100751
Fig. 1. (a) Calculated Fe-N phase diagram with N
2
isobars showing the phase stability of austenite and interface equilibrium conditions. (b) Experimental nitrogen
concentration distribution as determined with GD-OES. Light-optical microstructures of the as-annealed specimen (c) before and (d) after solution nitriding.
3.2. Phase stability of the SRB surface layer
X-ray diractograms from the surfaces of the SHT and HTSN speci-
mens before and after surface roller burnishing are given in Fig. 2 . The
more intense 220 austenite reection in for the HTSN specimen as com-
pared to the SHT specimen is attributed to additional grain growth in
nitrogen containing austenite (cf. Fig. 1 d). Evidently, such grain growth
promotes a slight change in texture favoring 110 oriented austenite
grains. The response of the SHT and HTSN surfaces towards SRB is very
dierent. Whilst the diractogram of the surface zone of the SHT + SRB
specimen shows mainly (b.c.c.) 𝛼′ -martensite after SRB, the HTSN + SRB
specimens only show austenite. Evidently, austenite is eectively sta-
bilized by HTSN, resulting in austenite that is fully resistant against
deformation-induced martensite formation upon severe deformation.
The line proles as measured in SHT and HTSN specimens are signi-
cantly broadened after SRB, which is attributed to a reduction in the co-
herently diracting domain size and/or an increase in the micro-strains
[26] .
To further clarify the evolution of the phase structure along with the
gradients of nitrogen concentration and strain in the deformed layer,
the phase composition was assessed with X-ray diraction at dierent
depths in the SRB treated specimens after successively removing sub-
layers by electropolishing (see Fig. 3 a and b). In the SHT + SRB spec-
imen austenite as well as b.c.c. 𝛼′ and h.c.p. 𝜀 martensite diraction
peaks are identied. The 𝛼′ -martensite peaks are present up to a rela-
tively large depth ( > 600 µm) and become more intense towards the
surface; 𝜀 -martensite is observed in the depth range 19 – 410 µm. For
the HTSN specimen deformed at the same feed depth, only austenite
diraction peaks are detected to a depth of 405 µm, while martensite
does appear beyond this depth, where the nitrogen content is very low
(cf. Fig. 1 b). Quantitative estimation of phase fractions was carried out
using the direct comparison method [27] , according to ASTM standard
E795 [28] . In the quantication procedure, the integrated intensity of
the 200 and 220 reections of austenite, the 220 and 211 reections of
𝛼′ -martensite, and of the 004 reection of 𝜀 -martensite were considered.
Fig. 3 c presents the corresponding depth distributions of the marten-
site volume fraction on these two treated specimens. For the SHT + SRB
specimen a sharp drop from 78 % to 46 % martensite occurs within
the rst ~20 µm from the surface, followed by a gradual reduction of
the martensite fraction. These ndings are consistent with those pre-
sented in Refs. [29 , 30] , where the amount of martensite was found to
increase with the degree of deformation and a reduction of the grain size
of AISI 304 stainless steel during cold working at room temperature. For
the HTSN + SRB specimen, beyond 405 µm depth, the small amount of
martensite detected ( < 2%) is comparable to the content determined in
the SHT + SRB specimen at this depth ( Fig. 3 c). The signicant dierence
in the volume fraction of deformation-induced martensite between the
SHT + SRB and HTSN + SRB specimens indicates that for the present con-
ditions, the combination of nitrogen content and degree of deformation
is sucient to suppress martensite formation up to a depth of ~ 400 µm.
The in-depth volume fraction of martensite in the SHT + SRB specimen
in Fig. 3 d, shows a considerable amount of 𝜀 -martensite with respect to
𝛼′ -martensite, especially in a range of medium degree of deformation.
For a higher feed depth of 500 µm, X-ray diractograms at dier-
ent depths for the HTSN + SRB specimen contain tiny diraction peaks
of 𝜀 -martensite at certain depths ( Fig. 3 b). Evidently, austenite is not
B. Wang, C. Hong and G. Winther et al. Materialia 12 (2020) 100751
Fig. 2. X-ray diractograms showing the phases present in the surface-adjacent
zone of the solution treated/nitrided specimens before and after SRB.
suciently stable under this more severe deformation condition. These
results suggest that 𝜀 -martensite is an intermediate phase during the
deformation-induced transformation 𝛾 → 𝜀 → 𝛼′ in the HTSN + SRB spec-
imen, and indicate that there exists a critical combination of nitrogen
content and degree of deformation where austenite is stable against
deformation-induced martensite formation.
3.3. Microstructure characterization
3.3.1. Overall observations in the deformed zone
Cross-sectional optical micrographs of the SHT + SRB and HTSN + SRB
specimens are given in Fig. 4 a-c, demonstrating dierent features
of plastic deformation and microstructure renement, corresponding
to dierent pre-treatment/SRB conditions. Irrespective of the pre-
treatment and straining conditions, in all materials SRB produced a sub-
stantial number of thin bands. A dense outermost zone with a thickness
of several tens of microns can be observed in all treated specimens, sug-
gesting the formation of ultrane/nano-crystalline grains (see below). In
good accordance with other severely surface deformed austenitic stain-
less steels [31–34] , the plastic deformation and strain induced by SRB
depend on the depth below the surface. The surface has experienced
the highest plastic strain, consisting of compression, as imposed by the
feed, and shear, through frictional contact with the roller. Plastic strain
gradually decreases with depth and the shear component vanishes, as
suggested by absence of grain boundaries that bend towards the for-
ward roller direction. In the following the deformation response of the
three treated specimens is described in separate sections.
For the SHT specimen, SRB with a feed depth of 300 µm has induced
a deformed zone (below the dense surface zone) that consists of multiple
thin straight bands the density of which increases towards the surface
( Fig. 4 a). Abundant band intersections develop in the upper part of the
deformed zone, while groups of parallel bands and individual bands are
visible deeper in the deformed zone.
Under the same SRB conditions, relatively few bands and hardly
any band intersections are observed in the surface-adjacent region of
the HTSN specimen ( Fig. 4 b), where a relatively high nitrogen content
is present. More pronounced formation of thin bands is observed
deeper (~ 430 µm) in the deformed zone, where the nitrogen content
is low, and appears to extend slightly deeper than for the SHT + SRB
specimen ( Fig. 4 d). The thickness of the deformed zones reects
the transferability of the plastic deformation, which depends on the
dissipation of the mechanical work exerted on the specimen during
SRB. Based on the processing characteristic of SRB, i.e. “repetitive
mechanical loading ”, the deformed microstructure developed at an
earlier stage would always participate in the energy dissipation during
the later deformation. The deeper deformed zone exists in HTSN + SRB
specimen reveals that deformation-induced twin/dislocation possesses
a higher transferability of plastic deformation than for martensite. This
may be attributed to the relatively high hardness of martensite (see
below the hardness results in Fig. 8 ), which limits the transfer of plastic
deformation to deeper regions. Similar deformation morphologies are
identied for a feed depth of 500 µm; the thicknesses of the dense
surface-adjacent zone and the entire deformed zone are larger than for
a 300 µm feed depth ( Fig. 4 c and d).
It is noted that the microstructure in the surface-adjacent region of
the deformed SHT specimen appears dark after etching for ~10 s, while
the deformed surface layer on the HTSN-300 specimen remains bright,
even for a relatively long etching time of 24 s (the HTSN-500 specimen
is slightly electrochemically attacked in this region). This dierence in
electrochemical response is explained from the largely martensitic sur-
face zone for the SHT + SRB and the nitrogen-containing austenitic sur-
face zone for the HTSN + SRB specimens.
3.3.2. Deformed microstructure on SHT + SRB specimen
The microstructure evolution along the deformation direction in the
SHT + SRB specimen is revealed by ECCI ( Fig. 5 a). Similar to that given
in Fig. 4 a, the thin bands are present in dierent congurations depend-
ing on depth: close to the surface, bands of a few micrometers thick are
intersected abundantly, while deeper in the zone sets of parallel bands
occur within the grains. The microstructure in the ultrane surface zone
as observed with light-optical microscopy ( Fig. 4 a) is not shown, because
this part was largely dissolved during electropolishing of the cross sec-
tion, consistent with its dark appearance in Fig. 4 a.
A combination of ECCI and EBSD was used in three selected re-
gions, corresponding to dierent depths in the deformed zone. In re-
gion I, the thin bands have a width of 1~2 µm and consist of a mix-
ture of 𝜀 -martensite and 𝛼′ -martensite. In particular, along a single 𝜀
lath there is a high density of 𝛼′ -martensite, some of which are wider
than the 𝜀 laths. These observations suggest that 𝛼′ -martensite has nu-
cleated inside the 𝜀 laths and thereafter has grown into the austenite
matrix. Similar structures were observed in Refs. [35–37] , where the
growth of 𝛼′ -martensite was reported to be conned by several adjacent
bands. At a depth of ~ 270 µm (region II) both individual bands and
groups of bands are present. Both 𝛼′ - and 𝜀 -martensite form in the indi-
vidual bands and are separated from the neighboring phases by a sharp
boundary, leading to serrated shape distribution, whereas exclusively
𝛼′ -martensite is observed in the groups of bands ( Fig. 5 e-g). It is most
likely that this 𝛼′ -martensite results from a direct 𝛾 → 𝛼′ transforma-
tion without 𝜀 -martensite as a pre-cursor for 𝛼′ -martensite, which has
also been observed previously [38 , 39] . At a depth of ~ 450 µm (re-
gion III), where a relatively low degree of deformation was applied, a
small number of intersected bands is found. At the intersections small
units of 𝛼′ -martensite have developed while the bands themselves con-
sist of 𝜀 -martensite ( Fig. 5 h-j). The development of 𝛼′ -martensite at band
intersections is consistent with the commonly accepted nucleation the-
ory stating that the intersections of 𝜀 -martensite or microscopic shear
bands including faults and twins, are preferred nucleation sites for 𝛼′ -
martensite [30 , 40 , 41] .
In concert with Fig. 3 d, the applied plastic deformation gradi-
ent in strain/strain rate resulted in uctuating fractions of marten-
site (both 𝛼′ - and 𝜀 -martensite) with depth in the SHT + SRB speci-
men, without playing a signicant role in altering the deformation
mechanism ( Fig. 5 ).
