Bulk nanocrystalline stainless steel fabricated by equalchannel angular pressing

C.X. Huanga)

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academyof Sciences, Shenyang 110016, People’s Republic of China

Y.L. Gao and G. YangCentral Iron and Steel Research Institute, Beijing 100081, People’s Republic of China

S.D. Wu,b) G.Y. Li, and S.X. LiShenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academyof 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 ultralowcarbon stainless steel by means of equal channel angular pressing at room temperature.Transmission electron microscopy (TEM) and high-resolution TEM investigationsindicated that two types of nanostructures were formed: nanocrystalline strain-inducedmartensite (body-centered cubic structure) with a mean grain size of 74 nm andnanocrystalline austenite (face-centered cubic structure) with a size of 31 nmcharacterized by dense deformation twins. The results about the formation of fullynanocrystalline grain structures in stainless steel suggested that a low stacking faultenergy is exceptionally profitable for producing nanocrystalline materials by equalchannel angular pressing.

I. INTRODUCTION

Nanocrystalline (nc) metals and alloys, conventionallydefined as polycrystals with grain sizes less than 100 nm,have exhibited superior mechanical properties, such asexcellent superplasticity and high strength.1–3 During thelast two decades, many severe plastic deformation (SPD)methods were developed for producing nanostructuredmaterials.4,5 Of these SPD methods, the equal channelangular pressing (ECAP) technique is the most promis-ing due to its ability to process bulk materials in threedimensions.4 However, the grain sizes obtained by ECAPfor many materials are actually outside the nc regime,i.e., on the order of several hundred nanometers (com-monly referred to as the ultrafine-grained range). Table Ipresents 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 asCu,6 Al alloys,7 Fe,8 low carbon steel,9 and Ti.10 For hardmaterials, such as Ti–6Al–4V12 and W,13 even at hightemperature, only subgrains with a size of submicrometerare formed. It has been shown that these materials wererefined via dislocation-controlled grain subdivision

mechanism and cannot be refined down to nanometers byECAP at room temperature (RT), especially for thosematerials with relatively high stacking fault energy(SFE).6–9 However, for materials with low SFE, the plas-tic deformation mode may change from dislocation slipto deformation twinning, and this is very important forgrain refinement. The mechanism of deformation twinsleading to both grain subdivision and a martensite trans-formation was identified in AISI 304 stainless steel dur-ing surface mechanical attrition treatment, and at thesame time, a nanostructured surface with grain sizes ofseveral tens of nanometers was obtained finally.14 Re-cently, Yapici et al.15 pressed 316L stainless steel at hightemperatures by ECAP and found deformation twinningin this material even at 800 °C but failed to producenanostructures. Therefore, in this work, we chose an ul-tralow carbon austenite stainless steel with a low SFE(∼20 mJ/m2)16 as the starting material and demonstratethat truly nc grain structures were achieved by means ofECAP at RT.

II. EXPERIMENTAL

The material used in this investigation was an ultralowcarbon austenite stainless steel with a composition, inweight 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 initialrod with a diameter of 8 mm and a length of 45 mm was

Address all correspondence to these authors.a)e-mail: cx_huang@imr.ac.cnb)e-mail: shdwu@imr.ac.cnDOI: 10.1557/JMR.2006.0214

J. Mater. Res., Vol. 21, No. 7, Jul 2006 © 2006 Materials Research Society 1687

annealed at 1150 °C for 2 h, which resulted in a grainsize in the range of 200–400 �m.

The ECAP procedure was performed using a die fab-ricated from tool steel (AISI M4-like) with two channelsintersecting at inner angle of 90°and outer angle of 30°.4

The rod coated with a MoS2 lubricant was pressed for8 passes (route Bc; i.e., the rod was rotated round thelongitudinal axis by 90° counterclockwise before eachpass4) 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 wasused to determine the phase constitution. The microstruc-ture observations were performed on a JEM-2000FX IItransmission electron microscope (TEM, operating at200 kV) and a Tecnai G2 S-Twin F30 high-resolutionTEM (HRTEM, operating at 300 kV, FEI Company).The thin foils for TEM and HRTEM observations werecut from the center of the pressed rod perpendicular tothe longitudinal axis of the rod, mechanically ground toabout 40 �m, and finally thinned by twin-jet polishingmethod (in a solution of 10% perchloric acid and ethanolat RT).

III. RESULTS AND DISCUSSION

A. X-ray diffraction analysis

Figure 1 shows x-ray diffraction (XRD) profiles of theas-received and the as-ECAP’ed samples. It can be foundthat the microstructure of the as-received sample is com-posed only of austenite, and the ECAP’ed one consists oflarge fraction of �� martensite. Quantitative XRD meas-urements indicated that the volume fraction of �� mar-tensite was ∼83%. Apparently, a strain-induced marten-site transformation took place during the ECAP treat-ment. As indicated by Shin et al.,17 shear deformationimposed by ECAP is the most effective method for in-troducing strain-induced martensite transformation com-pared with uniaxial compression and tensile deformation.

B. Nanostructures characterized by TEMand HRTEM

Figures 2(a) and 2(b) present typical bright-field andcorresponding dark-field TEM images, respectively, of

the ECAP’ed sample. It is obvious that the microstruc-tures are characterized by both equiaxed and elongatedgrains with sizes mostly on the nanometer scale. Thecorresponding selected-area diffraction pattern (SADP)taken from the region with a diameter of 1 �m shows thatall these nanograins are only martensites [body-centeredcubic (bcc) structure] with random crystallographic ori-entations. The histogram of grain size distribution [asshown in Fig. 2(c)] obtained from both bright- and dark-field TEM images (more than 500 grains were measured)shows a broad grain size distribution of 10–200 nm, and78% of the grains are smaller than 100 nm. The meangrain size determined by normal logarithmic distributionis approximately 74 nm.

Figure 3(a) shows another type of grain structuresformed in the same ECAP’ed sample. It can be seen thatthe grains are smaller and more uniform than thoseshown in Fig. 2(a). The corresponding SAD pattern in-dicates that they are austenite [face-centered cubic (fcc)structure] with random crystallographic orientations.Moreover, most of these grains contain two flat inter-faces parallel to each other [some of them are accentu-ated in the white circles in Fig. 3(a)]. The width andorientation of these planar defects vary from grain tograin, which is better illustrated in the dark-field TEMimage [Fig. 3(b); some of them are marked with whitecircles]. HRTEM observations (see next paragraph) in-dicate that they are deformation twins. Grain size meas-urements from both bright- and dark-field images show anarrow size distribution of 5–90 nm and the mean grainsize is about 31 nm [as shown in Fig. 3(c)]. The forma-tion of these fcc nanograins in low SFE stainless steelprobably resulted from different grain refinement processcompared with that of cubic materials with medium-highSFEs. For instance, in Inconel 600 alloy (fcc structurewith low SFE), the formation of nanograins during

TABLE I. Stable grain size (D) obtained by ECAP and the processingconditions of several materials.

Materials D (nm) T (°C) Passes Reference

Cu 270 RT 10 6Al–3 wt% Mg 270 RT 8 7Fe 235 RT 8 8Low-carbon steel 200 RT 4 9Ti 200 400 8 10Mg–Li 500 130 4 11Ti6Al4V 600 700 8 12W 1000 1000 3 13

FIG. 1. XRD profiles of the as-received and the ECAPed samples.

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surface mechanical attrition treatment involved the inter-action of microtwins and dislocations.18 They found thata large amount of deformation twins were first formed ininitial large grains, and subsequently, high-density dislo-cation arrays were induced inside the twin-matrix lamel-lae with a thickness of several tens of nanometers. These

dislocation arrays were finally evolved into high-anglegrain boundaries with further straining, subdividing thelamellae into nanograins.

Figure 4(a) is a typical HRTEM image viewed fromzone axis of [011]. Multiple deformation twins are de-tected on both (11̄1) plane (indicated by black arrows)

FIG. 2. TEM micrographs showing the martensite nanograins: (a)bright-field image and (b) dark-field image. (c) Grain size distributionwas determined from TEM observations. The inset of (a) shows thecorresponding SAD pattern.

FIG. 3. TEM micrographs showing the austenite nanograins: (a)bright-field image and (b) dark-field image. (c) Grain size distributionwas determined from TEM observations. The inset of (a) shows thecorresponding SAD pattern.

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and (111̄) plane (indicated by white arrows). A closeexamination 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 andstacking faults do not transect the entire grain, but ter-minate in grain interior in the middle parts of the imageas indicated by the two black arrows. It is obvious thatthese microtwins and stacking faults were nucleated atthe grain boundary and grew into the grain interior via

partial dislocation (Shockley type with Burgers vectors1/6[112]) emission from the grain boundary. Such atwinning mechanism in nanocrystallites has been pre-dicted by molecular dynamics simulations19 in nc Al andexperimentally evidenced in nc Al and Cu.20–22 Theubiquitousness of deformation twins implies that twin-ning via partial dislocation emission from grain boundaryis the primary deformation mode of nc austenite steel.Furthermore, the formation of deformation microtwins inturn refines the nanograins into much finer nanometer-sized blocks.

The above TEM investigations show that, evidently,fully nc grain structures have been formed in ultralowcarbon stainless steel by ECAP. Compared with that ofthe materials with medium to high SFEs, the grain re-finement mechanism of the material with low SFE mayshow different features. In previous work, Hansen andhis coworkers systematically studied the microstructuralevolution in cold-rolled fcc metals with medium to highSFEs, such as Cu and Al.23,24 They concluded that thegrain subdivision involves various dislocation activities.Severe plastic deformation generates high-density dislo-cations arranged into various configurations dependingon the nature of materials, such as the geometrically nec-essary boundary, incidental dislocation boundary, anddense dislocation wall.23,24 With increasing strain, someof these dislocation boundaries evolve into high-anglegrain boundaries that refine the original large grains intofiner grains.24 By means of ECAP, the stable grainsize that can be obtained is ∼1 �m for pure Al (SFE,166 mJ/m2),6,16,25 ∼450 nm for Al–1%Mg (SFE,110mJ/m2),7,16 and ∼270 nm for pure Cu (SFE,78 mJ/m2).6,16 The failures to reduce grain size down tothe nanometer scale are mainly due to the fast dynamicrecovery at RT that opposes the accumulation of dislo-cations and grain boundaries.6 This effect is the same forbcc materials with high SFE, such as Fe and low carbonsteel (see Table I). To produce nc grains, more rigorousdeformation conditions in addition to large strain are alsorequired. For example, Wang et al.26 rolled Cu to ex-tremely large strain at liquid-nitrogen temperature andobtained completely nc grains. They suggested that cryo-genic rolling led to the high accumulation of dislocationsthat facilitated dynamic recrystallization through copiousnucleation and growth, resulting in truly nc grain forma-tion. Another example is the formation of nanostructuredsurface layers by means of surface mechanical attritiontreatment. By peening the surface layer of Fe plate atvery high strain rate (103 to 104 s−1), a thin nanostruc-tured layer with grain size of 10–20 nm was formed at thetop surface.27 It is known that dislocation activity is sen-sitive to temperature and strain rate. Both low tempera-ture and high strain rate depress the dislocation activities,and therefore, higher dislocation density and finer grainsare expected to be obtained at very low temperature and/

FIG. 4. (a) HRTEM image of a nc grain viewed from [011] zone axis,showing multiple deformation twins and stacking faults, indicated byblack and white arrows. (b) High magnification of the rectangular in(a). The {111}-plane forming twin relationship is highlighted by whitelines. Many stacking faults are emitted from grain boundary and ter-minated in grain interior, as indicated by black arrows.

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or high strain rate straining. The function of low SFE issimilar to that of low temperature and high strain rate.The annihilation and rearrangement of dislocations intolower energy configurations during recovery are knownto be achieved by glide, climb, and cross-slip of dislo-cations.28 The climb of extended edge dislocations in fccmetals is controlled by vacancy evaporation at extendedjogs, and under these conditions the rate of climb is givenby28

� =cD

kT� F�2 ,

where c is the jog concentration, D is the coefficient ofdiffusion (D � e−Q/RT), k is the Boltzmann constant, T istemperature, F is the driving force of dislocation climb,and � is the SFE of a metal. Clearly, the recovery rate issubstantially affected by SFE (� � �2). Decreasing SFEof the metals decreases the velocity of dislocations re-markably and therefore suppresses the rate of recovery.High densities of dislocations result in the formation oflow-angle subgrain boundaries on a scale of nanometersto form nanocrystallites. These subgrain boundaries in-crease their misorientations with further straining, result-ing in the formation of nanograins with random orienta-tions.

Fcc structural materials with low SFEs tend to deformvia the mode of deformation twinning, but not disloca-tion slip. A large amount of deformation twins with finethickness, possibly in the submicrometer and nanometerregimes, may be formed under extremely high strain dur-ing the beginning several passes. These will result ingrain subdivision in a regime finer than those of thematerials with high SFE subdivided by dislocationboundaries. Furthermore, for stainless steel, martensitetransformation may occur within twin-matrix intersec-tions on a much finer scale, which is instrumental inrefining grains into the nanometer regime. Systematicinvestigations of the strain-induced phase transformationand grain refinement process of stainless steel underECAP deformation are in progress.

IV. SUMMARY

In summary, bulk fully nc grain structures have beensuccessfully achieved in low-carbon stainless steel bymeans of ECAP at RT. Two separate types of nc grainsare formed: strain-induced martensite with a mean grainsize of ∼74 nm and austenite with a size of ∼31 nm. It isconcluded that a low stacking fault energy is especiallyfavorable for the formation of nanocystalline grains byECAP at RT.

ACKNOWLEDGMENTS

The authors are thankful for the financial support fromNatural Science Foundation of China under Grant Nos.

50171072, 50371090, and 50471082. The authors ex-press their appreciation to Professor Z.F. Zhang for hisvaluable suggestions.

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