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.
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A surface layer with a depth-dependent microstructure was produced on a ferritic steel (P92) plate by means of surface mechanical
attrition treatment (SMAT). The austenitization processes of ferrite and carbides in the surface layers with different average grain sizes
were investigated by means of in situ X-ray diffraction, differential scanning calorimetry and transmission electron microscopy. Exper-
especially in steels, are a central topic in physical metal-
at high temperatures [6,7,12]. The case is more complicated
in high-alloy steels because the carbides in the initial
steel is relatively sluggish while growth of austenite is rapid,
The on-heating evolutions, such as dislocation recovery,
grain growth and carbide precipitation, as well as the resul-
tant mechanical properties, have been studied in nano-
structured or ultrafine-grained ferritic steels [20–22].
However, to the authors’ knowledge, there is no study on
the austenitization process of nanostructured ferrite
⇑ Corresponding authors. Tel.: +86 24 2397 1508; fax: +86 24 2399 8660.
E-mail addresses: zbwang@imr.ac.cn (Z.B. Wang), lu@imr.ac.cn (K.
Lu).
Available online at www.sciencedirect.com
Acta Materialia 59 (2011) 3710–37
lurgy because of a combination of fundamental scientific
interests and technological importance [1–4]. Among phase
transformations in steels, the on-heating formation of face-
centered cubic (fcc) austenite (c) from body-centered cubic
(bcc) ferrite (a) matrix has been studied extensively and
some investigations have dealt with the effects of alloying
elements and initial microstructures on the austenitization
process [5–13]. Experimental works on low-alloy steels
have indicated that cementite particles provide nucleation
sites for austenite, and austenite formation is very rapid
and the dissolution of carbides occurs in the austenite but
not in the ferrite. By studying the austenite formation
and the carbide dissolution in Fe–8.2Cr–C alloys with dif-
ferent C concentrations, Shtansky et al. [11] noticed that
the mechanisms of austenite nucleation and growth depend
on the composition, starting microstructure and austenitiz-
ing temperature.
Nanostructured materials have attracted intensive inter-
est for several decades, due to their novel properties origi-
nating from a large volume fraction of interfaces [15–19].
imental results showed that the onset temperature of the austenitization process decreases gradually with decreasing sizes of ferrite grains
and carbide particles, being �120 K lower in the top SMAT surface layer compared with that in the original sample. In addition, the two-
step austenitization process in the surface layers becomes a one-step one when the mean size of carbide particles is smaller than 20 nm.
The effects of microstructure refinement on the accelerated austenitization processes were discussed in terms of thermodynamic and
kinetic.
� 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Nanostructured; Surface mechanical attrition treatment; Ferritic steel; Grain size effects; Austenitization process
1. Introduction
Solid-state phase transformations in metallic materials,
structure are thermodynamically more stable than cement-
ite [14]. For example, Lenel and Honeycombe [10] observed
that nucleation of austenite in an Fe–10Cr–0.2C (wt.%)
Grain size effects on the
in a nanostructu
L.M. Wang, Z.B
Shenyang National Laboratory for Materials Science, Institute of Metal Rese
Received 23 December 2010; received in revised
Abstract
1359-6454/$36.00 � 2011 Acta Materialia Inc. Published by Elsevier Ltd. All
doi:10.1016/j.actamat.2011.03.006
ustenitization process
ed ferritic steel
Wang ⇑, K. Lu ⇑
h, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China
rm 25 February 2011; accepted 5 March 2011
www.elsevier.com/locate/actamat
19
rights reserved.
ater
matrix. This might be related to the fact that significant
growth of nanosized grains may occur upon heating prior
to the phase transformation temperature being reached.
At ambient temperature, c-Fe has been experimentally
observed in nanostructured Fe when the grain size is small
enough [23,24], and it was suggested to be thermodynami-
cally stable by calculating the Gibbs free energies of inter-
faces in nanostructured c and a grains [25]. Therefore,
notable grain size effects on the austenitization behaviors
in ferritic steels might be expected.
By means of surface mechanical attrition treatment
(SMAT), surface layers with a gradient grain size distribu-
tion (ranging from nanometers, submicrons to microns)
have been synthesized on various metallic materials
[26–34] as a result of gradient variations of applied strains
and strain rates with depth from the treated surface. Previ-
ous studies showed that the nanostructures with enhanced
Cr diffusivity in the SMAT surface layers of low-carbon
steel [35] and H13 steel [36] are effectively stabilized by
the formation of fine dispersive Cr compound particles
during chromizing at 873 K, resulting in the growth of
much thicker chromized surface layers than that on the
coarse-grained samples after chromizing treatments at tem-
peratures as high as 1323 K. This means that the gradient
microstructure of the SMAT surface layer on steels is sta-
ble at elevated temperatures with dispersive precipitates.
Such a kind of gradient nanostructured surface layer pro-
vides a unique opportunity to study the grain size effect
on austenitization behavior on the nanometer scale in
one sample.
In this work, a nanostructured surface layer with a
depth-dependent microstructure is synthesized on a com-
mercial ferritic steel plate by means of SMAT. Thermal sta-
bility and austenitization process in the surface layer are
characterized with respect to the microstructure.
2. Experimental
2.1. Sample preparation
The studied ferritic steel (P92) was supplied by Wyman
Gordon Forgings Inc., with the chemical composition (in
wt.%) of 0.11 C, 8.75 Cr, 0.40 Mo, 1.75 W, 0.18 V, 0.05
Nb, 0.05 N and balance Fe. The initial material was in
an austenitized and tempered condition (1 h at 1343 K fol-
lowed by 2.5 h at 1048 K). The plate sample
(100 � 50 � 4.0 mm3 in size) of the as-received steel was
submitted to SMAT, the set-up and procedure of which
have been described previously [26–28]. In brief, a large
number of hardened steel balls of 8 mm diameter were
placed at the bottom of a cylinder-shaped chamber and
vibrated at a high frequency by a generator. The sample
to be treated was fixed at the upper side of the chamber
and impacted by flying balls repeatedly and multidirection-
ally. Because the sample surface was plastically deformed
L.M. Wang et al. / Acta M
with high strains and high strain rates, grains in the surface
layer were effectively refined, and a depth-dependent
microstructure was consequently generated. In this work,
the sample was treated for 60 min in vacuum at ambient
temperature at a vibrating frequency of 50 Hz. No detect-
able contamination was introduced into the surface layer
during the SMAT process.
2.2. Microstructure characterization
The microstructure evolutions with depth in both the as-
SMAT and the annealed-SMAT samples were character-
ized using a JEM-2010 transmission electron microscope
(TEM) operated at a voltage of 200 kV. TEM foils of the
topmost layers were cut by the electro-spark discharge
technique, then mechanically polished, dimpled and finally
ion-milled from the untreated side. In addition, TEM foils
of different subsurface layers were cut, mechanically pol-
ished, dimpled and finally electropolished from the
untreated side of the SMAT samples after the removal of
surface layers of different thicknesses. The electropolishing
was carried out at 253 K with an electrolyte of 5 vol.% per-
chloric acid and 95 vol.% alcohol. A short-period milling
process at a low angle (at 4–5� for �10 min) from both
sides was applied to clean the foil surfaces before the
TEM observation. The grain/cell sizes were averaged from
a few hundred grains/cells selected randomly from TEM
images. In the case that not enough grains could be
counted in a TEM image at the greater depths, an FEI
Nova-nano scanning electron microscope (SEM) with a
lower magnification was applied to observe the microstruc-
ture at the corresponding region of a cross-sectional
sample.
The phase constitution of the original sample and the
SMAT surface layer at room temperature was identified
by X-ray diffraction (XRD) analysis using a Rigaku
D/max 2400 X-ray diffractometer (7.5 kW) with Cu Ka
radiation, with a step size of 0.02�.
2.3. Phase transformation measurements
The surface layers at different depths of the SMAT sam-
ple were cut by electro-spark discharge technique and then
mechanically polished from the untreated side to �10 lm
in thickness. Thermal analyses of the prepared foil samples
(�15 mg in weight) were conducted on a Netzsch differen-
tial scanning calorimeter (DSC 404C). The experiments
were carried out from room temperature to 1373 K at a
heating rate of 20 K min�1, in a flowing Ar atmosphere
with a gas flow rate of 50 ml min�1. The temperature was
calibrated by the melting points of pure In, Au and Ni.
In addition, the baseline in the region of each peak was
constructed by a polynomial (or an apparent transforma-
tion function) through the tangents at the left and right
sides of each peak at the evaluation limits [37]. The DSC
curves shown in the present work are the subtraction
results of the measured curves to the baselines. According
ialia 59 (2011) 3710–3719 3711
to the standards of the International Confederation for
Thermal Analysis, the onset temperature of each reaction
ter
(To) was determined as the intersecting point of the base-
line with the tangential line from the point with the maxi-
mum slope of the DSC curve at the left side of the peak.
In order to understand the phase transformation process
of the SMAT sample, an in situ XRD experiment was con-
ducted using a Brueker D8 Discover XRD (12 kW)
equipped with a high-temperature attachment. The temper-
ature was calibrated by the melting point of pure Al. The
measured sample was heated to preset temperatures at a
rate of 60 K min�1 and held for 1 min before collecting
the XRD profiles from 41.6 to 46�, with a step size of
0.02� and a scanning rate of 4� min�1. The sample temper-
ature was monitored by using a PtRh–Pt thermocouple, of
which the accuracy is ±2 K. All analyses were carried out
in vacuum (2 � 10�3 Pa).
3. Results and discussion
3.1. Microstructure of the SMAT sample
The initial microstructure of the sample before SMAT is
shown in Fig. 1a. A martensitic structure is observed with
numerous rod-like precipitates at lath/subgrain boundaries
in the austenitized and tempered ferritic steel. The width of
the laths is typically �540 nm and the length is �13 lm.
Subgrains are formed within laths during tempering to
reduce the density of dislocations formed by austenitizing
[38]. The precipitates are determined to be mostly of the
type M23C6 (M = Cr, Fe) according to the corresponding
selected area electron diffraction (SAED) pattern (see the
inset in Fig. 1a), with average dimensions of �80 nm along
the short axis and �160 nm along the long axis. The total
amount of precipitates is estimated to be �3 vol.% in the
original sample. In addition to M23C6 precipitates, various
much finer M(C, N) precipitates (�16 nm) were also
detected in the same material with the same heat treatment
by observing extraction double replicas [38,39].
Clear evidence of microstructure refining induced by
plastic deformation is observed in the SMAT surface layer
of �100 lm thickness. As shown in Fig. 1b, a high density
of dislocations are formed within the martensite (or ferrite;
the same in this work) laths and various dense dislocation
walls (DDWs) develop mostly parallel or perpendicular to
the lath directions. These are subdivided into smaller
grains/cells at a depth of �60 lm from the treated surface.
The grain/cell size is �300 nm along the short axis and
�600 nm along the long axis, respectively. Meanwhile, pre-
cipitates are also slightly refined to sizes of �60 nm along
the short axis and �120 nm along the long axis.
Further TEM observations indicate that more and more
dislocations and DDWs develop in the martensite laths and
the resulted grain/cell size decreases gradually with
decreasing depth from the treated surface, due to increas-
ing strain and strain rate. Moreover, the refined ferrite
grains/cells appear to be equiaxed when their size is below
3712 L.M. Wang et al. / Acta Ma
�35 nm (at a depth of <30 lm). As for the precipitates
(M23C6), the sizes along both axes also decrease gradually
with decreasing depth, and equiaxed particles are formed
when the size is below �20 nm (at a depth of <40 lm).
Due to the very high strain and strain rate (102–103 s�1
[27,28]), extremely fine equiaxed ferrite grains with random
crystallographic orientations are formed in the top surface
layer, as revealed by TEM observations and the corre-
sponding SAED pattern in Fig. 1c and d. The average size
of ferrite grains is estimated to be �8 nm from a number of
dark-field images taken from the (1 1 0) diffraction of a-Fe.
In addition, significant decreases in both the size and vol-
ume fraction of precipitates are observed in the top surface
layer. The mean size of M23C6 particles is refined to be
�4 nm and the volume fraction is reduced to be �1%,
according to the estimated results from a number of
dark-field images taken from the (3 1 1) diffraction of
M23C6 phase (see Fig. 1e).
Microstructure observations of the SMAT surface layer
indicate that both ferrite and precipitate grains in the pres-
ent ferritic steel are refined by dislocation activities, which
are similar to those observed previously in AISI 52100 steel
[29] and AISI H13 steel [40] during SMAT. In brief, the
refinement mechanism of ferrite grains involves formation
of DDWs and dislocation tangles in both the original
grains and the refined cells (under further straining), trans-
formation of these microstructures into subboundaries
with small misorientations, and evolution of subboundaries
to highly misoriented grain boundaries [28,29]. When fer-
rite grains are refined to a certain size, plastic deformation
occurs in precipitate particles due to grain refinement
strengthening of the ferrite matrix. Therefore, the particles
are also progressively refined into smaller particles and/or
dissolved into the ferrite matrix, as suggested by significant
decrements in both the size and volume fraction of M23C6
particles in the SMAT surface layer.
The measurement results of average sizes of ferrite
grains/cells and M23C6 particles are summarized in Fig. 2
as a function of depth from the SMAT surface. It is clear
that the top 40 lm surface layer is nanostructured (with
grain sizes below 100 nm) and the average sizes of both
phases increase with increasing depth in the top 100 lm
surface layer. Both ferrite and M23C6 grains appear to be
equiaxed in the top surface layer, but anisotropic morphol-
ogies are observed in the subsurface layer at a depth larger
than �30 lm.
3.2. Austenitization process of the SMAT ferritic steel
3.2.1. Phase transformation in the top surface layer during
heating
A typical DSC curve of the top surface layer (�10 lm in
thickness) of the SMAT sample is given in Fig. 3, in com-
parison with the DSC curve of the original sample without
SMAT. An exothermic peak (CP) between 800 and 870 K
is observed on the DSC curve of the SMAT sample, while
no such a peak appears on the curve of the original sample.
ialia 59 (2011) 3710–3719
A series of TEM observations and XRD analyses of micro-
structure evolutions of ferrite and precipitates across this
Fig. 1. Typical bright-field TEM images of (a) the original P92 steel sample and (b) at a depth of �60 lm of the SMAT sample. The inset in (a) shows the
corresponding SAED pattern of M23C6. (c) A typical bright-field TEM image of the top SMAT layer. (d and e) Dark-field TEM images of ferrite grains
and M23C6 particles, taken from diffractions of (1 1 0)a and (3 1 1)M23C6, respectively, as indicated on the SAED pattern (inset in (c)).
0 50 100 150
1
10
100
1000
Depth from surface (μm)
Si
ze
(n
m
)
Ferrite grain/cell (short axis)
Ferrite grain/cell (long axis)
M23C6 particle (short axis)
M23C6 particle (long axis)
Fig. 2. Variations of average grain/cell sizes of ferrite and M23C6 with
depth from the treated surface of the SMAT sample.
800 900 1000 1100 1200 1300
1019 K
AI
AIFM
FM
CP
H
ea
t F
lo
w
Temperature (K)
Original0.1 W/g
1122 K
AII
SMAT
Fig. 3. DSC curves of the original sample and the top surface layer of the
SMAT sample, at a heating rate of 20 K min�1.
L.M. Wang et al. / Acta Materialia 59 (2011) 3710–3719 3713
(b)γ(111) (a) γ(111)
F
nd
le w
3714 L.M. Wang et al. / Acta Mater
temperature range in an ongoing study (Wang et al., in
preparation) confirmed that this peak is mostly induced
by the significantly coarsening process of ferrite grains, as
well as the reforming and/or coarsening processes of pre-
cipitates from the broken and/or dissolved precursors dur-
ing SMAT. The endothermic peaks FM on the DSC curves
of both the SMAT sample and the original are confirmed
to be induced by the magnetic transition from a ferromag-
netic state to a paramagnetic state. No phase transition is
detected across the peaks and they are completely revers-
ible during DSC scans when the samples are heated to tem-
peratures below the second (AI) peaks and then cooled
down. Moreover, the peak positions of FM on the mea-
sured DSC curves agree very well with the reported Currie
point (�1018 K) of a reduced activation ferritic–martens-
42 44 46 48 42 44 46 48
In
te
ns
ity
(a
.u.
) SMAT
α(110)
1073 K
1023 K
1153 K
963 K
973 K
2 Theta (deg.)
1163 K
Original
α(110)
1173 K
1123 K
1263 K
1083 K
1093 K
1273 K
Fig. 4. In situ XRD profiles of (a) the topmost layer of the SMAT sample a
ratio F (derived by Eq. (1)) of the top SMAT layer and the original samp
itic steel with a nominal composition (wt.%) of 9Cr–
0.09C–1W [8].
Comparing the DSC curves of both samples after the
ferromagnetic transition, one can see that the onset temper-
ature of the endothermic peak AI on the curve of the
SMAT sample (1019þ26�2 K
1) is significantly decreased rela-
tive to that of the original sample (1122 ± 2 K). In addi-
tion, the broad endothermic peak AII observed in the
original sample disappears in the SMAT sample.
In situ XRD analyses were carried out on the SMAT
and the original samples from room temperature to
1373 K, as shown in Fig. 4a and b, respectively. A decrease
in the diffraction intensity of ferrite accompanied by an
increase in the diffraction intensity of austenite is observed
across the temperature intervals of peak AI for the SMAT
sample and of peaks AI and AII for the original sample.
While a quantitative determination of the volume fractions
1 Due to the overlap of FM and AI peaks, the baseline within this
temperature range could not be accurately determined. The uncertainty of
the onset temperature of AI in the SMAT sample is determined by using
baselines before the peak FM and after the peak AI, respectively.
of a and c phases during the in situ XRD measurements is
difficult, a factor F is defined to indicate the variation of
relative amount of c with temperature (see Fig. 4c), i.e.
F ¼ Icð111Þ
Icð111Þ þ Iað110Þ ð1Þ
where Ic(111) and Ia(110) are the diffraction intensities of c
(1 1 1) and a (1 1 0) peaks, respectively. Austenite is first
detected in the SMAT sample at the temperature of
973 K, its volume fraction increases with increasing tem-
perature and the sample is completely composed of austen-
ite at 1163 K. A gradual decrease in the fraction of a
accompanies the increase in c within this temperature
range. That is to say, the austenitization process in the
SMAT ferritic steel starts at
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