Bulk nanocrystalline stainless steel fabricated by equal channel

Bulk nanocrystalline stainless steel fabricated by equal
channel angular pressing
C.X. Huanga)
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy
of Sciences, Shenyang 110016, People’s Republic of China
Y.L. Gao and G. Yang
Central Iron and Steel Research Institute, Beijing 100081, People’s Republic of China
S.D. Wu,b) G.Y. Li, and S.X. Li
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy
of Sciences, Shenyang 110016, People’s Republic of China
(Received 26 September 2005; accepted 5 January 2006)
Bulk fully nanocrystalline grain structures were successfully obtained in ultralow
carbon stainless steel by means of equal channel angular pressing at room temperature.
Transmission electron microscopy (TEM) and high-resolution TEM investigations
indicated that two types of nanostructures were formed: nanocrystalline strain-induced
martensite (body-centered cubic structure) with a mean grain size of 74 nm and
nanocrystalline austenite (face-centered cubic structure) with a size of 31 nm
characterized by dense deformation twins. The results about the formation of fully
nanocrystalline grain structures in stainless steel suggested that a low stacking fault
energy is exceptionally profitable for producing nanocrystalline materials by equal
channel angular pressing.
I. INTRODUCTION
Nanocrystalline (nc) metals and alloys, conventionally
defined as polycrystals with grain sizes less than 100 nm,
have exhibited superior mechanical properties, such as
excellent superplasticity and high strength.1–3 During the
last two decades, many severe plastic deformation (SPD)
methods were developed for producing nanostructured
materials.4,5 Of these SPD methods, the equal channel
angular pressing (ECAP) technique is the most promising due to its ability to process bulk materials in three
dimensions.4 However, the grain sizes obtained by ECAP
for many materials are actually outside the nc regime,
i.e., on the order of several hundred nanometers (commonly referred to as the ultrafine-grained range). Table I
presents several typical materials processed by ECAP.
As shown, the finest grain size, typically of the order of
∼200 nm, is obtained in relatively soft materials, such as
Cu,6 Al alloys,7 Fe,8 low carbon steel,9 and Ti.10 For hard
materials, such as Ti–6Al–4V12 and W,13 even at high
temperature, only subgrains with a size of submicrometer
are formed. It has been shown that these materials were
refined via dislocation-controlled grain subdivision
Address all correspondence to these authors.
a)
e-mail: [email protected]
b)
e-mail: [email protected]
DOI: 10.1557/JMR.2006.0214
J. Mater. Res., Vol. 21, No. 7, Jul 2006
mechanism and cannot be refined down to nanometers by
ECAP at room temperature (RT), especially for those
materials with relatively high stacking fault energy
(SFE).6–9 However, for materials with low SFE, the plastic deformation mode may change from dislocation slip
to deformation twinning, and this is very important for
grain refinement. The mechanism of deformation twins
leading to both grain subdivision and a martensite transformation was identified in AISI 304 stainless steel during surface mechanical attrition treatment, and at the
same time, a nanostructured surface with grain sizes of
several tens of nanometers was obtained finally.14 Recently, Yapici et al.15 pressed 316L stainless steel at high
temperatures by ECAP and found deformation twinning
in this material even at 800 °C but failed to produce
nanostructures. Therefore, in this work, we chose an ultralow carbon austenite stainless steel with a low SFE
(∼20 mJ/m2)16 as the starting material and demonstrate
that truly nc grain structures were achieved by means of
ECAP at RT.
II. EXPERIMENTAL
The material used in this investigation was an ultralow
carbon austenite stainless steel with a composition, in
weight percent, of 0.007 C, 18.46 Cr, 11.82 Ni, 1.61 Si,
0.008 S, 0.018P. 0.29 Mn, and the balance Fe. The initial
rod with a diameter of 8 mm and a length of 45 mm was
© 2006 Materials Research Society
1687
C.X. Huang et al.: Bulk nanocrystalline stainless steel fabricated by equal channel angular pressing
TABLE I. Stable grain size (D) obtained by ECAP and the processing
conditions of several materials.
Materials
D (nm)
T (°C)
Passes
Reference
Cu
Al–3 wt% Mg
Fe
Low-carbon steel
Ti
Mg–Li
Ti6Al4V
W
270
270
235
200
200
500
600
1000
RT
RT
RT
RT
400
130
700
1000
10
8
8
4
8
4
8
3
6
7
8
9
10
11
12
13
annealed at 1150 °C for 2 h, which resulted in a grain
size in the range of 200–400 ␮m.
The ECAP procedure was performed using a die fabricated from tool steel (AISI M4-like) with two channels
intersecting at inner angle of 90°and outer angle of 30°.4
The rod coated with a MoS2 lubricant was pressed for
8 passes (route Bc; i.e., the rod was rotated round the
longitudinal axis by 90° counterclockwise before each
pass4) at RT at a pressing speed of 9 mm/min.
A Rigaku D/max-2400 x-ray diffractometer (12 kW,
Rigaku Corporation, Japan) with Cu K␣ radiation was
used to determine the phase constitution. The microstructure observations were performed on a JEM-2000FX II
transmission electron microscope (TEM, operating at
200 kV) and a Tecnai G2 S-Twin F30 high-resolution
TEM (HRTEM, operating at 300 kV, FEI Company).
The thin foils for TEM and HRTEM observations were
cut from the center of the pressed rod perpendicular to
the longitudinal axis of the rod, mechanically ground to
about 40 ␮m, and finally thinned by twin-jet polishing
method (in a solution of 10% perchloric acid and ethanol
at RT).
III. RESULTS AND DISCUSSION
A. X-ray diffraction analysis
Figure 1 shows x-ray diffraction (XRD) profiles of the
as-received and the as-ECAP’ed samples. It can be found
that the microstructure of the as-received sample is composed only of austenite, and the ECAP’ed one consists of
large fraction of ␣⬘ martensite. Quantitative XRD measurements indicated that the volume fraction of ␣⬘ martensite was ∼83%. Apparently, a strain-induced martensite transformation took place during the ECAP treatment. As indicated by Shin et al.,17 shear deformation
imposed by ECAP is the most effective method for introducing strain-induced martensite transformation compared with uniaxial compression and tensile deformation.
B. Nanostructures characterized by TEM
and HRTEM
Figures 2(a) and 2(b) present typical bright-field and
corresponding dark-field TEM images, respectively, of
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FIG. 1. XRD profiles of the as-received and the ECAPed samples.
the ECAP’ed sample. It is obvious that the microstructures are characterized by both equiaxed and elongated
grains with sizes mostly on the nanometer scale. The
corresponding selected-area diffraction pattern (SADP)
taken from the region with a diameter of 1 ␮m shows that
all these nanograins are only martensites [body-centered
cubic (bcc) structure] with random crystallographic orientations. The histogram of grain size distribution [as
shown in Fig. 2(c)] obtained from both bright- and darkfield TEM images (more than 500 grains were measured)
shows a broad grain size distribution of 10–200 nm, and
78% of the grains are smaller than 100 nm. The mean
grain size determined by normal logarithmic distribution
is approximately 74 nm.
Figure 3(a) shows another type of grain structures
formed in the same ECAP’ed sample. It can be seen that
the grains are smaller and more uniform than those
shown in Fig. 2(a). The corresponding SAD pattern indicates that they are austenite [face-centered cubic (fcc)
structure] with random crystallographic orientations.
Moreover, most of these grains contain two flat interfaces parallel to each other [some of them are accentuated in the white circles in Fig. 3(a)]. The width and
orientation of these planar defects vary from grain to
grain, which is better illustrated in the dark-field TEM
image [Fig. 3(b); some of them are marked with white
circles]. HRTEM observations (see next paragraph) indicate that they are deformation twins. Grain size measurements from both bright- and dark-field images show a
narrow size distribution of 5–90 nm and the mean grain
size is about 31 nm [as shown in Fig. 3(c)]. The formation of these fcc nanograins in low SFE stainless steel
probably resulted from different grain refinement process
compared with that of cubic materials with medium-high
SFEs. For instance, in Inconel 600 alloy (fcc structure
with low SFE), the formation of nanograins during
J. Mater. Res., Vol. 21, No. 7, Jul 2006
C.X. Huang et al.: Bulk nanocrystalline stainless steel fabricated by equal channel angular pressing
FIG. 2. TEM micrographs showing the martensite nanograins: (a)
bright-field image and (b) dark-field image. (c) Grain size distribution
was determined from TEM observations. The inset of (a) shows the
corresponding SAD pattern.
FIG. 3. TEM micrographs showing the austenite nanograins: (a)
bright-field image and (b) dark-field image. (c) Grain size distribution
was determined from TEM observations. The inset of (a) shows the
corresponding SAD pattern.
surface mechanical attrition treatment involved the interaction of microtwins and dislocations.18 They found that
a large amount of deformation twins were first formed in
initial large grains, and subsequently, high-density dislocation arrays were induced inside the twin-matrix lamellae with a thickness of several tens of nanometers. These
dislocation arrays were finally evolved into high-angle
grain boundaries with further straining, subdividing the
lamellae into nanograins.
Figure 4(a) is a typical HRTEM image viewed from
zone axis of [011]. Multiple deformation twins are detected on both (11̄1) plane (indicated by black arrows)
J. Mater. Res., Vol. 21, No. 7, Jul 2006
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C.X. Huang et al.: Bulk nanocrystalline stainless steel fabricated by equal channel angular pressing
FIG. 4. (a) HRTEM image of a nc grain viewed from [011] zone axis,
showing multiple deformation twins and stacking faults, indicated by
black and white arrows. (b) High magnification of the rectangular in
(a). The {111}-plane forming twin relationship is highlighted by white
lines. Many stacking faults are emitted from grain boundary and terminated in grain interior, as indicated by black arrows.
and (111̄) plane (indicated by white arrows). A close
examination of the white rectangular area in Fig. 4(a)
shows a high density of microtwins and Stacking Faults
(SF) [Fig. 4(b)]. As shown, some of these microtwins and
stacking faults do not transect the entire grain, but terminate in grain interior in the middle parts of the image
as indicated by the two black arrows. It is obvious that
these microtwins and stacking faults were nucleated at
the grain boundary and grew into the grain interior via
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partial dislocation (Shockley type with Burgers vectors
1/6[112]) emission from the grain boundary. Such a
twinning mechanism in nanocrystallites has been predicted by molecular dynamics simulations19 in nc Al and
experimentally evidenced in nc Al and Cu.20–22 The
ubiquitousness of deformation twins implies that twinning via partial dislocation emission from grain boundary
is the primary deformation mode of nc austenite steel.
Furthermore, the formation of deformation microtwins in
turn refines the nanograins into much finer nanometersized blocks.
The above TEM investigations show that, evidently,
fully nc grain structures have been formed in ultralow
carbon stainless steel by ECAP. Compared with that of
the materials with medium to high SFEs, the grain refinement mechanism of the material with low SFE may
show different features. In previous work, Hansen and
his coworkers systematically studied the microstructural
evolution in cold-rolled fcc metals with medium to high
SFEs, such as Cu and Al.23,24 They concluded that the
grain subdivision involves various dislocation activities.
Severe plastic deformation generates high-density dislocations arranged into various configurations depending
on the nature of materials, such as the geometrically necessary boundary, incidental dislocation boundary, and
dense dislocation wall.23,24 With increasing strain, some
of these dislocation boundaries evolve into high-angle
grain boundaries that refine the original large grains into
finer grains.24 By means of ECAP, the stable grain
size that can be obtained is ∼1 ␮m for pure Al (SFE,
166 mJ/m 2 ), 6,16,25 ∼450 nm for Al–1%Mg (SFE,
110mJ/m 2 ), 7,16 and ∼270 nm for pure Cu (SFE,
78 mJ/m2).6,16 The failures to reduce grain size down to
the nanometer scale are mainly due to the fast dynamic
recovery at RT that opposes the accumulation of dislocations and grain boundaries.6 This effect is the same for
bcc materials with high SFE, such as Fe and low carbon
steel (see Table I). To produce nc grains, more rigorous
deformation conditions in addition to large strain are also
required. For example, Wang et al.26 rolled Cu to extremely large strain at liquid-nitrogen temperature and
obtained completely nc grains. They suggested that cryogenic rolling led to the high accumulation of dislocations
that facilitated dynamic recrystallization through copious
nucleation and growth, resulting in truly nc grain formation. Another example is the formation of nanostructured
surface layers by means of surface mechanical attrition
treatment. By peening the surface layer of Fe plate at
very high strain rate (103 to 104 s−1), a thin nanostructured layer with grain size of 10–20 nm was formed at the
top surface.27 It is known that dislocation activity is sensitive to temperature and strain rate. Both low temperature and high strain rate depress the dislocation activities,
and therefore, higher dislocation density and finer grains
are expected to be obtained at very low temperature and/
J. Mater. Res., Vol. 21, No. 7, Jul 2006
C.X. Huang et al.: Bulk nanocrystalline stainless steel fabricated by equal channel angular pressing
or high strain rate straining. The function of low SFE is
similar to that of low temperature and high strain rate.
The annihilation and rearrangement of dislocations into
lower energy configurations during recovery are known
to be achieved by glide, climb, and cross-slip of dislocations.28 The climb of extended edge dislocations in fcc
metals is controlled by vacancy evaporation at extended
jogs, and under these conditions the rate of climb is given
by28
␷=
cD
⭈ F␥2
kT
,
where c is the jog concentration, D is the coefficient of
diffusion (D ⬀ e−Q/RT), k is the Boltzmann constant, T is
temperature, F is the driving force of dislocation climb,
and ␥ is the SFE of a metal. Clearly, the recovery rate is
substantially affected by SFE (␷ ⬀ ␥2). Decreasing SFE
of the metals decreases the velocity of dislocations remarkably and therefore suppresses the rate of recovery.
High densities of dislocations result in the formation of
low-angle subgrain boundaries on a scale of nanometers
to form nanocrystallites. These subgrain boundaries increase their misorientations with further straining, resulting in the formation of nanograins with random orientations.
Fcc structural materials with low SFEs tend to deform
via the mode of deformation twinning, but not dislocation slip. A large amount of deformation twins with fine
thickness, possibly in the submicrometer and nanometer
regimes, may be formed under extremely high strain during the beginning several passes. These will result in
grain subdivision in a regime finer than those of the
materials with high SFE subdivided by dislocation
boundaries. Furthermore, for stainless steel, martensite
transformation may occur within twin-matrix intersections on a much finer scale, which is instrumental in
refining grains into the nanometer regime. Systematic
investigations of the strain-induced phase transformation
and grain refinement process of stainless steel under
ECAP deformation are in progress.
IV. SUMMARY
In summary, bulk fully nc grain structures have been
successfully achieved in low-carbon stainless steel by
means of ECAP at RT. Two separate types of nc grains
are formed: strain-induced martensite with a mean grain
size of ∼74 nm and austenite with a size of ∼31 nm. It is
concluded that a low stacking fault energy is especially
favorable for the formation of nanocystalline grains by
ECAP at RT.
ACKNOWLEDGMENTS
The authors are thankful for the financial support from
Natural Science Foundation of China under Grant Nos.
50171072, 50371090, and 50471082. The authors express their appreciation to Professor Z.F. Zhang for his
valuable suggestions.
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