《金属材料强韧化与组织调控》教学资源（参考文献）An investigation of surface nanocrystallization mechanism in Fe induced by surface mechanical attrition treatment

cta MATERIALIA Pergamon Acta Materialia 50(2002)4603-4616 www.actamat-journals.com An investigation of surface nanocrystallization mechanism in Fe induced by surface mechanical attrition treatment N.R.Tao",Z.B.Wang",W.P.Tong,M.L.Suia,J.Lu b,K.Lu a* .Shenyang National Laboratory for Materials Science,Institute of Metal Research.Chinese Academy of Sciences.72 Wenhua Road.Shenyang 110016.People's Republic of China LASMIS.University of Technology of Troyes.10000 Troyes.France Received 10 April 2002;received in revised form 4 June 2002;accepted 18 July 2002 Abstract By means of surface mechanical attrition (SMA),a nanostructured surface layer was formed on a pure Fe plate Microstructure features of various sections in the surface layer,from the strain-free matrix to the treated top surface, were systematically characterized by using X-ray diffraction(XRD)analysis,scanning electron microscopy(SEM)and transmission electron microscopy (TEM)observations.Based on the experimental observations,a grain refinement mechanism induced by plastic deformation during the SMA treatment in Fe was proposed.It involves formation of dense dislocation walls(DDWs)and dislocation tangles(DTs)in original grains and in the refined cells(under further straining)as well,transformation of DDWs and DTs into subboundaries with small misorientations separating individual cells or subgrains,and evolution of subboundaries to highly misoriented grain boundaries.Experimental evidences and analysis of the grain refinement mechanism indicate that high strains with a high strain rate are necessary for formation of nanocrystallites during plastic deformation of metals. 2002 Acta Materialia Inc.Published by Elsevier Science Ltd.All rights reserved. Keywords:Nanocrystalline materials;Iron;Grain refinement;Microstructure 1.Introduction bological properties [4].The workability of nc metals is much improved due to their enhanced Nanocrystalline (nc)materials,which are struc- superplasticity at lower temperatures compared to turally characterized by nanometer-sized grains the conventional polycrystals [5,6].These proper- with a large number of grain boundaries,have been ties and performance enable nc materials to be found to exhibit many novel properties relative to potentially very useful in developing new material their coarse-grained counterparts [1,2].For families with much enhanced properties for engin- example,most nc metals and alloys possess high eering applications and in upgrading the manufac- strength and hardness [3],as well as excellent tri- turing process of traditional engineering materials in industries. Nevertheless,research and development of bulk Corresponding author.Fax:+86-24-2399-8660. nc materials are hindered,to some extent,by vari- E-mail address:lu@imr.ac.cn (K.Lu). ous difficulties in synthesis techniques.Since the 1359-6454/02/$22.00 2002 Acta Materialia Inc.Published by Elsevier Science Ltd.All rights reserved. Pl:S1359-6454(02)00310-5

Acta Materialia 50 (2002) 4603–4616 www.actamat-journals.com An investigation of surface nanocrystallization mechanism in Fe induced by surface mechanical attrition treatment N.R. Tao a , Z.B. Wang a , W.P. Tong a , M.L. Sui a , J. Lu b , K. Lu a,∗ a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, People’s Republic of China b LASMIS, University of Technology of Troyes, 10000 Troyes, France Received 10 April 2002; received in revised form 4 June 2002; accepted 18 July 2002 Abstract By means of surface mechanical attrition (SMA), a nanostructured surface layer was formed on a pure Fe plate. Microstructure features of various sections in the surface layer, from the strain-free matrix to the treated top surface, were systematically characterized by using X-ray diffraction (XRD) analysis, scanning electron microscopy (SEM) and transmission electron microscopy (TEM) observations. Based on the experimental observations, a grain refinement mechanism induced by plastic deformation during the SMA treatment in Fe was proposed. It involves formation of dense dislocation walls (DDWs) and dislocation tangles (DTs) in original grains and in the refined cells (under further straining) as well, transformation of DDWs and DTs into subboundaries with small misorientations separating individual cells or subgrains, and evolution of subboundaries to highly misoriented grain boundaries. Experimental evidences and analysis of the grain refinement mechanism indicate that high strains with a high strain rate are necessary for formation of nanocrystallites during plastic deformation of metals.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Nanocrystalline materials; Iron; Grain refinement; Microstructure 1. Introduction Nanocrystalline (nc) materials, which are struc￾turally characterized by nanometer-sized grains with a large number of grain boundaries, have been found to exhibit many novel properties relative to their coarse-grained counterparts [1,2]. For example, most nc metals and alloys possess high strength and hardness [3], as well as excellent tri- ∗ Corresponding author. Fax: +86-24-2399-8660. E-mail address: lu@imr.ac.cn (K. Lu). 1359-6454/02/$22.00  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 13 59 -6454(02)00310-5 bological properties [4]. The workability of nc metals is much improved due to their enhanced superplasticity at lower temperatures compared to the conventional polycrystals [5,6]. These proper￾ties and performance enable nc materials to be potentially very useful in developing new material families with much enhanced properties for engin￾eering applications and in upgrading the manufac￾turing process of traditional engineering materials in industries. Nevertheless, research and development of bulk nc materials are hindered, to some extent, by vari￾ous difficulties in synthesis techniques. Since the

4604 N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 pioneering study of preparing bulk nc metals by approach to create localized plastic deformation, means of gas condensation and consolidation by resulting in grain refinement progressively down to Gleiter et al.,in the early 1980s,several processing the nanometer region in the surface layer of met- techniques have been developed to produce bulk nc allic materials [12-14].The SMA process,which materials,e.g.consolidation of ultrafine powders has been successfully applied in many material prepared by various kinds of techniques [7],crys- systems [15],has some unique advantages com- tallization of amorphous precursors [8],ball-mill- pared with the coating and deposition methods for ing and consolidation [9],severe plastic defor- SNC.For example,as there is no change in chemi- mation of bulk metals [10],and electrodeposition cal compositions of the nc surface layer and in the [11].However,due to the limitation of each of matrix,as well as a gradient variation in the grain these techniques,preparation of 'ideal'bulk nc dimension from nano-sized (in the top layer)to samples (free of contamination and porosity,bulk coarse-grains (matrix),bonding of the nc surface in size,uniform and small (of a few nanometers)in layer with matrix will not be a problem.In grain size)is still a challenge to material scientists. addition,many existing mechanical processing Meanwhile much effort has been concentrated on techniques are applicable for synthesizing nanos- improvement of these techniques for synthesizing tructures in the surface by modifying the pro- bulk nc materials in recent years,the techniques cessing parameters,such as shot peening,hammer for producing nanostructured surface layers are peening,surface rolling,laser shock processing,etc emerging quickly. [12].Therefore,this SNC approach has a great Most failures of materials occur on surfaces, potential in industrial applications.For further including fatigue fracture,fretting fatigue,wear development of the SMA technique,a clear under- and corrosion etc.,which are very sensitive to the standing of the underlying mechanism for forma- structure and properties of the material surface. tion of nanostructures during the treatment is Optimization of the surface structure and proper- necessary.The objective of this work is to reveal ties may effectively enhance the global behavior of the intrinsic mechanism for grain refinement dur- a material.As a result,the surface modification of ing the SMA treatment in a pure iron sample. engineering materials is found to process more and Plastic deformation induced grain refinement more industrial applications.With increasing evi- has been known for many years [10,16].Polycrys- dences of novel properties in nc materials,it is tals with submicro-sized grains were usually fabri- reasonable to propose to achieve surface modifi- cated via severe plastic deformation (SPD)of vari- cation by the generation of a nanostructured sur- ous metals and alloys by using equal-channel face layer so that the overall properties and angular pressing (ECAP)[10],high pressure tor- behavior of the material are significantly improved. sion [17]and cold rolling [18].By examining the This kind of surface modification,referred as sur- microstructure evolution in plastically deformed face nanocrystallization (SNC),will greatly metals,possible grain refinement mechanisms were enhance the surface properties without changing proposed,involving dislocation activities,forma- the chemical composition [12].It is also a flexible tion of subgrain boundaries and grain boundaries approach that makes it possible to meet specific [18,19].The refinement process of coarse grains structure/property requirements on surface of upon plastic deformation,in principle,depends on samples. many intrinsic and extrinsic factors,such as struc- Nanostructured surface layers can be produced ture and stacking fault energy of the material,the by means of various existing coating and depo- intensity of strains and strain rates,deformation sition techniques such as PVD,CVD and plasma temperature,and so on.Formation of nano-sized processing.Alternatively,SNC can be realized by grains was realized in intensive mechanical means of grain refinement into the nanometer attrition processes (such as ball-milling)in metals regime in the surface layer of a bulk material.Our and alloys [20-22],in which much larger strains previous investigations demonstrated that surface and strain rates were applied relative to the other mechanical attrition (SMA)is an effective SPD processes.In the milling process,fracture and

4604 N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 pioneering study of preparing bulk nc metals by means of gas condensation and consolidation by Gleiter et al., in the early 1980s, several processing techniques have been developed to produce bulk nc materials, e.g. consolidation of ultrafine powders prepared by various kinds of techniques [7], crys￾tallization of amorphous precursors [8], ball-mill￾ing and consolidation [9], severe plastic defor￾mation of bulk metals [10], and electrodeposition [11]. However, due to the limitation of each of these techniques, preparation of ‘ideal’ bulk nc samples (free of contamination and porosity, bulk in size, uniform and small (of a few nanometers) in grain size) is still a challenge to material scientists. Meanwhile much effort has been concentrated on improvement of these techniques for synthesizing bulk nc materials in recent years, the techniques for producing nanostructured surface layers are emerging quickly. Most failures of materials occur on surfaces, including fatigue fracture, fretting fatigue, wear and corrosion etc., which are very sensitive to the structure and properties of the material surface. Optimization of the surface structure and proper￾ties may effectively enhance the global behavior of a material. As a result, the surface modification of engineering materials is found to process more and more industrial applications. With increasing evi￾dences of novel properties in nc materials, it is reasonable to propose to achieve surface modifi- cation by the generation of a nanostructured sur￾face layer so that the overall properties and behavior of the material are significantly improved. This kind of surface modification, referred as sur￾face nanocrystallization (SNC), will greatly enhance the surface properties without changing the chemical composition [12]. It is also a flexible approach that makes it possible to meet specific structure/property requirements on surface of samples. Nanostructured surface layers can be produced by means of various existing coating and depo￾sition techniques such as PVD, CVD and plasma processing. Alternatively, SNC can be realized by means of grain refinement into the nanometer regime in the surface layer of a bulk material. Our previous investigations demonstrated that surface mechanical attrition (SMA) is an effective approach to create localized plastic deformation, resulting in grain refinement progressively down to the nanometer region in the surface layer of met￾allic materials [12–14]. The SMA process, which has been successfully applied in many material systems [15], has some unique advantages com￾pared with the coating and deposition methods for SNC. For example, as there is no change in chemi￾cal compositions of the nc surface layer and in the matrix, as well as a gradient variation in the grain dimension from nano-sized (in the top layer) to coarse-grains (matrix), bonding of the nc surface layer with matrix will not be a problem. In addition, many existing mechanical processing techniques are applicable for synthesizing nanos￾tructures in the surface by modifying the pro￾cessing parameters, such as shot peening, hammer peening, surface rolling, laser shock processing, etc [12]. Therefore, this SNC approach has a great potential in industrial applications. For further development of the SMA technique, a clear under￾standing of the underlying mechanism for forma￾tion of nanostructures during the treatment is necessary. The objective of this work is to reveal the intrinsic mechanism for grain refinement dur￾ing the SMA treatment in a pure iron sample. Plastic deformation induced grain refinement has been known for many years [10,16]. Polycrys￾tals with submicro-sized grains were usually fabri￾cated via severe plastic deformation (SPD) of vari￾ous metals and alloys by using equal-channel angular pressing (ECAP) [10], high pressure tor￾sion [17] and cold rolling [18]. By examining the microstructure evolution in plastically deformed metals, possible grain refinement mechanisms were proposed, involving dislocation activities, forma￾tion of subgrain boundaries and grain boundaries [18,19]. The refinement process of coarse grains upon plastic deformation, in principle, depends on many intrinsic and extrinsic factors, such as struc￾ture and stacking fault energy of the material, the intensity of strains and strain rates, deformation temperature, and so on. Formation of nano-sized grains was realized in intensive mechanical attrition processes (such as ball-milling) in metals and alloys [20–22], in which much larger strains and strain rates were applied relative to the other SPD processes. In the milling process, fracture and

N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 4605 cold-welding of metal particles,as well as con- Vacuum tamination from the milling media are involved, Sample which make it difficult to identify the dominating mechanism for refining the grains into the nano- meter scales.Therefore,up to now,a clear scenery of formation of nano-sized crystallites from coarse polycrystals by plastic deformation is still lacking. In the SMA treatment,different microstructures Vibration can be obtained within the deformed surface layer generator along the depth from the treated surface to the strain-free matrix,i.e.from nano-sized grains to (a) (b) submicro-sized and micro-sized crystallites.This Fig.1. Schematic illustrations of the SMA treatment set-up gradient structure results from a gradient change in (a)and the localized plastic deformation in the surface layer by applied strains and strain rates along the depth, the impacting of the shot (b). from very large (top surface layer)to zero(strain- free matrix),in the treated surface layer.Obvi- the high vibration frequency of the system(50 Hz ously,the SMA treatment provides a unique in this work,which can be varied up to 20 kHz opportunity to investigate the grain refinement [12,131),the sample surface to be treated was mechanism by examining the microstructure fea- peened by a large number of shots in a short period tures at different depths in the deformed surface of time.Each peening of the ball to the surface will layer,in which the deformation-induced evidences result in plastic deformation in the surface layer of for grain refinement with different strains and the treated sample (as shown in Fig.1(b)).As a strain rates are maintained. consequence,repeated multidirectional peening at high strain rates onto the sample surface leads to severe plastic deformation in the surface layer.In 2.Experimental this work,the samples were treated in vacuum for 60 min at room temperature.After the SMA treat- 2.1.Sample ment,the sample surface is smooth (with compara- ble roughness as the original polished sample) An iron plate(70×l00×l00mm3 in size)with Positron annihilation spectroscopy experiments a purity of 99.95 wt%,was subjected to the SMA revealed that the surface layer of the SMA treated treatment in order to achieve a nc surface layer. sample is free of porosity. Before SMA treatments,the plate surfaces were polished with silicon carbide papers and then 2.3.Microstructural characterization annealed in vacuum at 1223 K for 120 min for diminishing the effect of mechanical processing X-ray diffraction (XRD)analysis of the surface and obtaining homogeneous coarse grains.The layer in the SMA treated Fe sample was carried grain size of the annealed Fe plate is on average out on a Rigaku D/max 2400 X-ray diffractometer 100-150um (12 kW),with Cu Ko radiation (wavelengths AKa 1.54056A and AKo2 1.54439A were 2.2.SMA treatment [15] reflected by a graphite crystal using the (0002) reflection).Small angular steps of 20 =0.02 were Fig.1(a)shows a schematic illustration of the taken to measure the intensity of each Bragg dif- SMA treatment set-up used in the present work. fraction peak.The counting time of 20 s was used Stainless steel balls (shots)of 8 mm in diameter to exactly measure the width of diffraction peak in were placed at the bottom of a cylinder-shaped the step-scanning mode.The average grain size vacuum chamber that was vibrated by a generator, and mean microstrain were calculated from line with which the shots were resonated.Because of broadening of bcc Fe (110),(200),(211),(220)

N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 4605 cold-welding of metal particles, as well as con￾tamination from the milling media are involved, which make it difficult to identify the dominating mechanism for refining the grains into the nano￾meter scales. Therefore, up to now, a clear scenery of formation of nano-sized crystallites from coarse polycrystals by plastic deformation is still lacking. In the SMA treatment, different microstructures can be obtained within the deformed surface layer along the depth from the treated surface to the strain-free matrix, i.e. from nano-sized grains to submicro-sized and micro-sized crystallites. This gradient structure results from a gradient change in applied strains and strain rates along the depth, from very large (top surface layer) to zero (strain￾free matrix), in the treated surface layer. Obvi￾ously, the SMA treatment provides a unique opportunity to investigate the grain refinement mechanism by examining the microstructure fea￾tures at different depths in the deformed surface layer, in which the deformation-induced evidences for grain refinement with different strains and strain rates are maintained. 2. Experimental 2.1. Sample An iron plate (70 × 100 × 100 mm3 in size) with a purity of 99.95 wt%, was subjected to the SMA treatment in order to achieve a nc surface layer. Before SMA treatments, the plate surfaces were polished with silicon carbide papers and then annealed in vacuum at 1223 K for 120 min for diminishing the effect of mechanical processing and obtaining homogeneous coarse grains. The grain size of the annealed Fe plate is on average 100–150 µm. 2.2. SMA treatment [15] Fig. 1(a) shows a schematic illustration of the SMA treatment set-up used in the present work. Stainless steel balls (shots) of 8 mm in diameter were placed at the bottom of a cylinder-shaped vacuum chamber that was vibrated by a generator, with which the shots were resonated. Because of Fig. 1. Schematic illustrations of the SMA treatment set-up (a) and the localized plastic deformation in the surface layer by the impacting of the shot (b). the high vibration frequency of the system (50 Hz in this work, which can be varied up to 20 kHz [12,13]), the sample surface to be treated was peened by a large number of shots in a short period of time. Each peening of the ball to the surface will result in plastic deformation in the surface layer of the treated sample (as shown in Fig. 1(b)). As a consequence, repeated multidirectional peening at high strain rates onto the sample surface leads to severe plastic deformation in the surface layer. In this work, the samples were treated in vacuum for 60 min at room temperature. After the SMA treat￾ment, the sample surface is smooth (with compara￾ble roughness as the original polished sample). Positron annihilation spectroscopy experiments revealed that the surface layer of the SMA treated sample is free of porosity. 2.3. Microstructural characterization X-ray diffraction (XRD) analysis of the surface layer in the SMA treated Fe sample was carried out on a Rigaku D/max 2400 X-ray diffractometer (12 kW), with Cu Kα radiation (wavelengths lKα1
1.54056A˚ and lKα2
1.54439A˚ were reflected by a graphite crystal using the (0002) reflection). Small angular steps of 2q
0.02° were taken to measure the intensity of each Bragg dif￾fraction peak. The counting time of 20 s was used to exactly measure the width of diffraction peak in the step-scanning mode. The average grain size and mean microstrain were calculated from line broadening of bcc Fe (110), (200), (211), (220)

4606 N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 (310)and(222)Bragg diffraction peaks,by using I reated. surface the Scherrer and Wilson method [23].By using repeated electrochemical etching,the treated sur- face layer was removed layer-by-layer,so that the microstructural evolution along the depth from the treated surface was determined by means of XRD analyses. Surface layer Cross-sectional observations of the treated Fe sample were preformed on a JSM-6301F scanning electron microscope (SEM).Microstructure fea- tures in the surface layer were characterized by using a Philip EM-420 transmission electron microscope (TEM,operating at a voltage of 120 Matrix kV)and a JEOL-2010 high-resolution electron microscope (HREM,at 200 kV),respectively. Plane-view and cross-sectional thin foils for TEM and HREM observations were prepared by means of cutting,grinding,dimpling and a final ion thin- ning at low temperatures. 50um 3.Results 3.1.Variation of grain size along depth Fig.2.A cross-sectional SEM observation of the SMA treated Fe sample for 60 min. Fig.2 is a cross-sectional SEM observation of the SMA treated Fe sample.Obviously,micro- peaks was seen owing to a grain refinement and structure morphology of the treated surface layer an increase in the atomic-level microstrain.The (of about 60 um thick)differs from that in the average grain size was calculated to be about 12 matrix.Severe plastic deformation evidences are nm and the mean microstrain is 0.12%.These seen in the surface layer,in which grain boundaries values are comparable to those for the ball-milled could not be identified as in the matrix.It is noted Fe powder sample reported [25].It implies a that the surface deformation layer thickness in the similar processing intensity between the SMA entire Fe plate surface is not uniform from place treatment(in the top surface layer)and the ball- to place (+20 um),indicating the heterogeneity of milling process. plastic deformation induced by the repeated peen- The XRD peak broadening becomes less evident ing.This may be attributed to the heterogeneous as the depth increases,indicating that grain sizes nature of the plastic deformation both within and gradually increase and/or the microstrain between grains [24]. decreases.When the grain size is larger than 100 XRD analyses were carried out for determining nm,the XRD analysis results become inaccurate the average grain (or cell)size and the microstrain due to the large uncertainty in measuring the peak in the surface layer.With consideration of the Cu broadening.TEM and SEM observations were Ko wavelength and its extinction depth in Fe,XRD employed to determine the grain size.Fig.3 shows patterns reflect the structure information from the a typical TEM plane-view observation of the top surface layer of about 5 um thick.By removal of surface layer(about 1 um deep)in the SMA treated the surface layer-by-layer,structure evolution Fe sample.It is clear that the microstructure is along the depth can be determined.For the top sur- characterized by ultrafine equiaxed grains(bcc Fe) face layer,evident broadening of Bragg diffraction with random crystallographic orientations,as indi-

4606 N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 (310) and (222) Bragg diffraction peaks, by using the Scherrer and Wilson method [23]. By using repeated electrochemical etching, the treated sur￾face layer was removed layer-by-layer, so that the microstructural evolution along the depth from the treated surface was determined by means of XRD analyses. Cross-sectional observations of the treated Fe sample were preformed on a JSM-6301F scanning electron microscope (SEM). Microstructure fea￾tures in the surface layer were characterized by using a Philip EM-420 transmission electron microscope (TEM, operating at a voltage of 120 kV) and a JEOL-2010 high-resolution electron microscope (HREM, at 200 kV), respectively. Plane-view and cross-sectional thin foils for TEM and HREM observations were prepared by means of cutting, grinding, dimpling and a final ion thin￾ning at low temperatures. 3. Results 3.1. Variation of grain size along depth Fig. 2 is a cross-sectional SEM observation of the SMA treated Fe sample. Obviously, micro￾structure morphology of the treated surface layer (of about 60 µm thick) differs from that in the matrix. Severe plastic deformation evidences are seen in the surface layer, in which grain boundaries could not be identified as in the matrix. It is noted that the surface deformation layer thickness in the entire Fe plate surface is not uniform from place to place (±20 µm), indicating the heterogeneity of plastic deformation induced by the repeated peen￾ing. This may be attributed to the heterogeneous nature of the plastic deformation both within and between grains [24]. XRD analyses were carried out for determining the average grain (or cell) size and the microstrain in the surface layer. With consideration of the Cu Kα wavelength and its extinction depth in Fe, XRD patterns reflect the structure information from the surface layer of about 5 µm thick. By removal of the surface layer-by-layer, structure evolution along the depth can be determined. For the top sur￾face layer, evident broadening of Bragg diffraction Fig. 2. A cross-sectional SEM observation of the SMA treated Fe sample for 60 min. peaks was seen owing to a grain refinement and an increase in the atomic-level microstrain. The average grain size was calculated to be about 12 nm and the mean microstrain is 0.12%. These values are comparable to those for the ball-milled Fe powder sample reported [25]. It implies a similar processing intensity between the SMA treatment (in the top surface layer) and the ball￾milling process. The XRD peak broadening becomes less evident as the depth increases, indicating that grain sizes gradually increase and/or the microstrain decreases. When the grain size is larger than 100 nm, the XRD analysis results become inaccurate due to the large uncertainty in measuring the peak broadening. TEM and SEM observations were employed to determine the grain size. Fig. 3 shows a typical TEM plane-view observation of the top surface layer (about 1 µm deep) in the SMA treated Fe sample. It is clear that the microstructure is characterized by ultrafine equiaxed grains (bcc Fe) with random crystallographic orientations, as indi-

N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 4607 (b Mean microstrain (XRD) Average grain/oell size [XRD) Grain/cell size (equiaxed,TEM) 0.15 Cell size (short axis.TEM) Ccl记(long axi成TEM) Coll sio (SEM) 0.10 10 102 0.05 microstrain 多 10 50 nm 50 nm 10 990.00 Fig.3.(a)Bright-field and (b)dark-field TEM images show- 10 20304050607080 ing planar microstructures of the top surface layer in the SMA Distance from surface (um) treated Fe sample (insert is a statistic distribution of grain size derived from the dark-field TEM images). Fig.4. Variations of the grain/cell size and the mean microstrain with the depth from the treated surface of the Fe sample determined by means of XRD analysis,TEM and SEM cated by the selected area electron diffraction observations.In TEM observations,sizes of the equiaxed cells (SAED)pattern.The histogram of grain size distri- and the lamellar cells(along short axis and long axis)are meas- ured. bution obtained from the dark-field images was characterized by a normal logarithmic distribution with a narrow size distribution (insert).The mean section (iii):micro-sized regime (40-60 um); grain size is approximately 7 nm,which is slightly section (iv):matrix with plastic deformation evi- smaller than the XRD result(12 nm).This might dences (60-110 um). be partially attributed to the fact that the XRD result averages the structure information of the top 3.2.Microstructure features surface layer of about 5 um thick.All diffraction rings in the SAED pattern were identified as bcc The experimental results clearly demonstrated Fe,and no other phases were detected. that nanostructures developed in the surface layer The microstrain determined from XRD analysis of the Fe plate during the treatment.In order to decreases significantly along the depth in the sur- understand the microstructure evolution process in face layer and gradually drops to zero at about 60 the surface layer during the treatment,a systematic um deep.From cross-sectional TEM and SEM investigation on the microstructure in the surface observations,we noticed evidences of dislocation layer at different stages of straining is needed.As activities induced by the attrition treatment in even the straining decreases from maximum at the top deeper matrix (up to 110 um deep),where no surface layer to zero in the matrix,the structure change in grain size and atomic-level microstrain evolution process during the SMA treatment may is detected.The measured average grain (or cell) be signed by the microstructure characteristics size and the mean microstrain as a function of (with different strains)at different depths from the depth in the SMA treated Fe sample by using dif- top surface to deep matrix.Therefore,detailed ferent analysis techniques (XRD,SEM,and TEM) cross-sectional TEM observations were performed are summarized in Fig.4.In terms of the grain (or on the treated Fe sample. cell)size,the SMA treated surface layer can be subdivided into four sections along depth from the 3.2.1.Dislocations,dense dislocation wall and top surface: dislocation tangle Fig.5 shows TEM observations in the defor- section (i):nanostructured regime (0-15 um); mation layer at low strains adjacent to the strain- section (ii):submicro-sized regime(15-40 um); free matrix(60-80 um deep from the top surface

N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 4607 Fig. 3. (a) Bright-field and (b) dark-field TEM images show￾ing planar microstructures of the top surface layer in the SMA treated Fe sample (insert is a statistic distribution of grain size derived from the dark-field TEM images). cated by the selected area electron diffraction (SAED) pattern. The histogram of grain size distri￾bution obtained from the dark-field images was characterized by a normal logarithmic distribution with a narrow size distribution (insert). The mean grain size is approximately 7 nm, which is slightly smaller than the XRD result (12 nm). This might be partially attributed to the fact that the XRD result averages the structure information of the top surface layer of about 5 µm thick. All diffraction rings in the SAED pattern were identified as bcc Fe, and no other phases were detected. The microstrain determined from XRD analysis decreases significantly along the depth in the sur￾face layer and gradually drops to zero at about 60 µm deep. From cross-sectional TEM and SEM observations, we noticed evidences of dislocation activities induced by the attrition treatment in even deeper matrix (up to 110 µm deep), where no change in grain size and atomic-level microstrain is detected. The measured average grain (or cell) size and the mean microstrain as a function of depth in the SMA treated Fe sample by using dif￾ferent analysis techniques (XRD, SEM, and TEM) are summarized in Fig. 4. In terms of the grain (or cell) size, the SMA treated surface layer can be subdivided into four sections along depth from the top surface: section (i): nanostructured regime (0–15 µm); section (ii): submicro-sized regime (15–40 µm); Fig. 4. Variations of the grain/cell size and the mean microstrain with the depth from the treated surface of the Fe sample determined by means of XRD analysis, TEM and SEM observations. In TEM observations, sizes of the equiaxed cells and the lamellar cells (along short axis and long axis) are meas￾ured. section (iii): micro-sized regime (40–60 µm); section (iv): matrix with plastic deformation evi￾dences (60–110 µm). 3.2. Microstructure features The experimental results clearly demonstrated that nanostructures developed in the surface layer of the Fe plate during the treatment. In order to understand the microstructure evolution process in the surface layer during the treatment, a systematic investigation on the microstructure in the surface layer at different stages of straining is needed. As the straining decreases from maximum at the top surface layer to zero in the matrix, the structure evolution process during the SMA treatment may be signed by the microstructure characteristics (with different strains) at different depths from the top surface to deep matrix. Therefore, detailed cross-sectional TEM observations were performed on the treated Fe sample. 3.2.1. Dislocations, dense dislocation wall and dislocation tangle Fig. 5 shows TEM observations in the defor￾mation layer at low strains adjacent to the strain￾free matrix (60–80 µm deep from the top surface

4608 N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 some regions the dislocation density is rather high.High-density dislocations arraying in tangles were observed,as in Fig.5(b).In tangles dislocations are randomly arranged without preferable sliding orientations. 3.2.2.Grain subdivision and subgrains As the depth decreases.deformation strains and strain rate increase.In the micro-sized regime (section(iii),40-60 um deep),most original grains (a) 1μm (b) 200nm are found to be subdivided into micro-sized cells (or 'blocks'),of which the shapes are either Fig.5.Cross-sectional TEM images in the matrix of the roughly equiaxed or lamellar in cross-sectional treated sample with plastic deformation evidences of DDWs TEM observations. ((a)numbers signifying misorientations across DDWs),and DTs((b)in which white arrows indicate location of high dislo- Fig.6 shows these regular-shaped cation density regimes). (parallelogram)cells separated by two sets of inter- secting DDWs in the {110)planes.These inter- secting DDWs seem to 'cut'the original coarse section (iv)),in which three typical deformation- grain into refined (micro-sized)blocks,across induced microstructure features were identified: which small misorientations (less than 1)are observed. 1.Dislocation lines (DLs):Homogeneously dis- tributed DLs are observed in Fe {110)planes and other sliding planes (such as (112)and (123)),depending upon the orientations of the grains.The density of DLs increases with a decrease of depth from the top surface. 2.Dense dislocation walls(DDWs):They are fre- quently seen inside some grains in this section. as shown in Fig.5(a).These DDWs (along Fe {110)planes)are parallel to each other separ- ated with a uniform spacing.It is noticed that the spacing between parallel DDWs varies (from a fraction of micron to a few microns) from grain to grain,depending upon grain orien- tations.DDWs are believed to result from dislo- cation accumulation and rearrangement for min- imizing the total energy state.It is worth noting that small misorientations across DDWs are detected,usually smaller than 1(as indicated in Fig.5(a)).Meanwhile,dislocation lines along the (112)planes are also visible,which inter- sect with DDWs in the {110)planes.Inter- secting DDWs are occasionally observed in 1μm some grains in which DDWs were developed simultaneously in different slip planes. Fig.6.A TEM image showing intersecting DDWs developed 3.Dislocation tangles (DTs):In some grains,the in two sets of(110)planes that cut the original grain into paral- dislocation distribution is not uniform and in lelogram cells

4608 N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 Fig. 5. Cross-sectional TEM images in the matrix of the treated sample with plastic deformation evidences of DDWs ((a) numbers signifying misorientations across DDWs), and DTs ((b) in which white arrows indicate location of high dislo￾cation density regimes). section (iv)), in which three typical deformation￾induced microstructure features were identified: 1. Dislocation lines (DLs): Homogeneously dis￾tributed DLs are observed in Fe {110} planes and other sliding planes (such as {112} and {123}), depending upon the orientations of the grains. The density of DLs increases with a decrease of depth from the top surface. 2. Dense dislocation walls (DDWs): They are fre￾quently seen inside some grains in this section, as shown in Fig. 5(a). These DDWs (along Fe {110} planes) are parallel to each other separ￾ated with a uniform spacing. It is noticed that the spacing between parallel DDWs varies (from a fraction of micron to a few microns) from grain to grain, depending upon grain orien￾tations. DDWs are believed to result from dislo￾cation accumulation and rearrangement for min￾imizing the total energy state. It is worth noting that small misorientations across DDWs are detected, usually smaller than 1° (as indicated in Fig. 5(a)). Meanwhile, dislocation lines along the {112} planes are also visible, which inter￾sect with DDWs in the {110} planes. Inter￾secting DDWs are occasionally observed in some grains in which DDWs were developed simultaneously in different slip planes. 3. Dislocation tangles (DTs): In some grains, the dislocation distribution is not uniform and in some regions the dislocation density is rather high. High-density dislocations arraying in tangles were observed, as in Fig. 5(b). In tangles dislocations are randomly arranged without preferable sliding orientations. 3.2.2. Grain subdivision and subgrains As the depth decreases, deformation strains and strain rate increase. In the micro-sized regime (section (iii), 40–60 µm deep), most original grains are found to be subdivided into micro-sized cells (or ‘blocks’), of which the shapes are either roughly equiaxed or lamellar in cross-sectional TEM observations. Fig. 6 shows these regular-shaped (parallelogram) cells separated by two sets of inter￾secting DDWs in the {110} planes. These inter￾secting DDWs seem to ‘cut’ the original coarse grain into refined (micro-sized) blocks, across which small misorientations (less than 1°) are observed. Fig. 6. A TEM image showing intersecting DDWs developed in two sets of {110} planes that cut the original grain into paral￾lelogram cells

N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 4609 Fig.7 shows lamellar-shaped cell structures with DDWs may become subboundaries with larger a width of about 400-1000 nm and a length of misorientations.As seen in Fig.7,intersecting about a few microns (section (ii)).Similar lamel- DDWs and subboundaries can form a parallelo- lar-shaped structures were also observed in tension gram structure.With the transformation of DDWs experiments of Fe [26,27].The lamellar cells are (along the {110)planes)into subboundaries with separated by evident boundaries'in the long axis large misorientations,roughly equiaxed cells sur- (along the (110)planes).These 'boundaries'are rounded by subboundaries are formed (see the cell obviously sharper and thinner than the DDWs(as labeled by 'X'). in Fig.6),and the misorientations across them are With an increase of strain at smaller depth,the much larger(usually a few degrees)than those for number of DDWs and subboundaries increases, DDWs.It is reasonable to believe these 'bound- and intersections of DDWs and subboundaries aries'(referred as 'subboundaries')are developed appear more frequently,so that smaller cells were from original DDWs by accumulation and annihil- observed.Fig.8 shows the cross-sectional micro- ation of more dislocations.By absorbing of more structures observed at the depth of about 24 um and more dislocations,the DDWs may transform from the top surface.The microstructures are into subboundaries (i.e.small-angle grain characterized by lamellar cells and subgrains,of boundaries)with an increased misorientations.In which the sizes are evidently smaller than those in Fig.7,some parts of DDWs along the long axis Fig.7.The size of lamellar cells along the long have not been transformed completely into the sub- axis ranges from about one to several micrometers, boundaries,as indicated.DDWs along the short and the short axis size is about 200-400 nm.Simi- axis of the cells are also observed to be along the lar to Fig.7,the submicro-sized lamellar cells are (110)planes,which intersect the subboundaries separated by either DDWs or subboundaries.Cells along the long axis.Misorientations across adjac- 2 and 3 along the short axis as well as cells 2 and ent DDWs were observed to be less than 1.Some 4 along the long axis are separated by DDWs, intersecting DDWs are tilted probably due to the which have not yet been transformed into subboun- mutual interaction with dislocations. daries.Subboundaries were found between cells The intersection of DDWs along the short axis 3/4,4/5 and 5/6.Development of some subbound- with subboundaries (long axis)results in further aries from the DDWs is seen in progress,e.g. refinement of the cells by 'cutting'the lamellar between cells 4/5 and 5/6,as indicated in Fig.8. cells into smaller cells.By accumulating more dis- It is noticed that these cells can be further div- locations with further straining,the cutting' ided into smaller lamellar or equiaxed substruc- tures via development of more DDWs and sub- boundaries.As one can see in cell 1 in Fig.8, b development of new DDWs parallel to the long axis is in progress (as indicated),which may eventually result in formation of subboundaries, leading to subdivision of the cell.Formation of DDWs in other directions inside a cell is also poss- ible,for example,in cell 6.DDWs normal to the long axis of cell 5/6 subboundary are formed, which may cut the cell into more cells with small length/thickness ratios.High density dislocations 1画 exist in some cells (as seen in cell 6),which can probably be subdivided by development of DT Fig.7.Cross-sectional TEM images at 37 um deep from the structures and subboundaries,as will be discussed top surface of the Fe sample.Submicro-sized lamellar cells sep- arated by DDWs (dotted lines)and subboundaries (solid lines) in the following section. are seen.Numbers indicate misorientations across adjacent Another typical microstructure observed in the cells. submicro-sized regime is characterized by equi-

N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 4609 Fig. 7 shows lamellar-shaped cell structures with a width of about 400–1000 nm and a length of about a few microns (section (ii)). Similar lamel￾lar-shaped structures were also observed in tension experiments of Fe [26,27]. The lamellar cells are separated by evident ‘boundaries’ in the long axis (along the {110} planes). These ‘boundaries’ are obviously sharper and thinner than the DDWs (as in Fig. 6), and the misorientations across them are much larger (usually a few degrees) than those for DDWs. It is reasonable to believe these ‘bound￾aries’ (referred as ‘subboundaries’) are developed from original DDWs by accumulation and annihil￾ation of more dislocations. By absorbing of more and more dislocations, the DDWs may transform into subboundaries (i.e. small-angle grain boundaries) with an increased misorientations. In Fig. 7, some parts of DDWs along the long axis have not been transformed completely into the sub￾boundaries, as indicated. DDWs along the short axis of the cells are also observed to be along the {110} planes, which intersect the subboundaries along the long axis. Misorientations across adjac￾ent DDWs were observed to be less than 1°. Some intersecting DDWs are tilted probably due to the mutual interaction with dislocations. The intersection of DDWs along the short axis with subboundaries (long axis) results in further refinement of the cells by ‘cutting’ the lamellar cells into smaller cells. By accumulating more dis￾locations with further straining, the ‘cutting’ Fig. 7. Cross-sectional TEM images at 37 µm deep from the top surface of the Fe sample. Submicro-sized lamellar cells sep￾arated by DDWs (dotted lines) and subboundaries (solid lines) are seen. Numbers indicate misorientations across adjacent cells. DDWs may become subboundaries with larger misorientations. As seen in Fig. 7, intersecting DDWs and subboundaries can form a parallelo￾gram structure. With the transformation of DDWs (along the {110} planes) into subboundaries with large misorientations, roughly equiaxed cells sur￾rounded by subboundaries are formed (see the cell labeled by ‘X’). With an increase of strain at smaller depth, the number of DDWs and subboundaries increases, and intersections of DDWs and subboundaries appear more frequently, so that smaller cells were observed. Fig. 8 shows the cross-sectional micro￾structures observed at the depth of about 24 µm from the top surface. The microstructures are characterized by lamellar cells and subgrains, of which the sizes are evidently smaller than those in Fig. 7. The size of lamellar cells along the long axis ranges from about one to several micrometers, and the short axis size is about 200–400 nm. Simi￾lar to Fig. 7, the submicro-sized lamellar cells are separated by either DDWs or subboundaries. Cells 2 and 3 along the short axis as well as cells 2 and 4 along the long axis are separated by DDWs, which have not yet been transformed into subboun￾daries. Subboundaries were found between cells 3/4, 4/5 and 5/6. Development of some subbound￾aries from the DDWs is seen in progress, e.g. between cells 4/5 and 5/6, as indicated in Fig. 8. It is noticed that these cells can be further div￾ided into smaller lamellar or equiaxed substruc￾tures via development of more DDWs and sub￾boundaries. As one can see in cell 1 in Fig. 8, development of new DDWs parallel to the long axis is in progress (as indicated), which may eventually result in formation of subboundaries, leading to subdivision of the cell. Formation of DDWs in other directions inside a cell is also poss￾ible, for example, in cell 6. DDWs normal to the long axis of cell 5/6 subboundary are formed, which may cut the cell into more cells with small length/thickness ratios. High density dislocations exist in some cells (as seen in cell 6), which can probably be subdivided by development of DT structures and subboundaries, as will be discussed in the following section. Another typical microstructure observed in the submicro-sized regime is characterized by equi-

4610 N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 500nm Fig.8.A cross-sectional TEM image in the submicro-sized section(24 um deep from the top surface)showing lamellar dislocation cells and subgrains with different dislocation configurations and boundaries (solid triangles for DDWs,open triangles for subboundar- ies,arrows for undeveloped subboundaries and letters I for location of high dislocation density regimes). axed cells (or subgrains).as shown in Fig.9.The (as evidenced by SAED pattern)are separated by size of cells is about 200-500 nm.and small mis- sharp conventional boundaries.Also,dislocations orientations were detected among these cells.As inside the grains are present,as in typical micro- in the lamellar-shaped cells,the equiaxed cells are structure developed by severe plastic deformation also separated by DDWs,DTs,or subboundaries, [281 as can be clearly seen in Fig.9.It is noticed that the subboundaries of the equiaxed cells have not 3.2.3.Nanostructures specific orientations (in the lamellar cells,the sub- The deformation strain and strain rate are drasti- boundaries are normally parallel to the slip planes). cally increased in the top surface layer.Nano-sized One may speculate that these subboundaries are microstructures were observed in the regime of developed from the high-density DTs in which no from 15 um deep to the top surface.Close to the preferable orientation exists.Development of DTs top surface,nano-sized equiaxed nanocrystallites into subboundaries can be seen in Fig.9,as indi- are seen,and lamellar-shaped nanograins are cated.The wide 'band'of DTs (indicated by solid present at the bottom of this section (adjacent to arrows)seem to form a new subboundary.Inside the submicro-sized section).Figs.11 and 12 show some cells one can see evident contrasts of a high typical cross-sectional observations of the micro- density of lattice dislocations,which may involve structure in this section.In Fig.11(a),lamellar in further refinement of the cells with increasing grains of 10-50 nm thick and 50-100 nm long are strains.Eventually,those subboundaries transform observed.It is worth noting that inside the inner into conventional grain boundaries with large mis- of these lamellar nanocrystallites,there exist even orientations by accumulating more dislocations, smaller equiaxed nano-grains.In Fig.11(b),one and orientations of the grains become random.Fig. can see roughly equiaxed nanocrystallites with 10 shows equiaxed grains with sizes of about 300 sizes of about 30 nm.The SAED pattern indicates nm observed at the depth of about 28 um from the small misorientations among these nanocrystal- top surface.These grains with random orientations lites.Evidently,these equiaxed nanograins were

4610 N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 Fig. 8. A cross-sectional TEM image in the submicro-sized section (24 µm deep from the top surface) showing lamellar dislocation cells and subgrains with different dislocation configurations and boundaries (solid triangles for DDWs, open triangles for subboundar￾ies, arrows for undeveloped subboundaries and letters I for location of high dislocation density regimes). axed cells (or subgrains), as shown in Fig. 9. The size of cells is about 200–500 nm, and small mis￾orientations were detected among these cells. As in the lamellar-shaped cells, the equiaxed cells are also separated by DDWs, DTs, or subboundaries, as can be clearly seen in Fig. 9. It is noticed that the subboundaries of the equiaxed cells have not specific orientations (in the lamellar cells, the sub￾boundaries are normally parallel to the slip planes). One may speculate that these subboundaries are developed from the high-density DTs in which no preferable orientation exists. Development of DTs into subboundaries can be seen in Fig. 9, as indi￾cated. The wide ‘band’ of DTs (indicated by solid arrows) seem to form a new subboundary. Inside some cells one can see evident contrasts of a high density of lattice dislocations, which may involve in further refinement of the cells with increasing strains. Eventually, those subboundaries transform into conventional grain boundaries with large mis￾orientations by accumulating more dislocations, and orientations of the grains become random. Fig. 10 shows equiaxed grains with sizes of about 300 nm observed at the depth of about 28 µm from the top surface. These grains with random orientations (as evidenced by SAED pattern) are separated by sharp conventional boundaries. Also, dislocations inside the grains are present, as in typical micro￾structure developed by severe plastic deformation [28]. 3.2.3. Nanostructures The deformation strain and strain rate are drasti￾cally increased in the top surface layer. Nano-sized microstructures were observed in the regime of from 15 µm deep to the top surface. Close to the top surface, nano-sized equiaxed nanocrystallites are seen, and lamellar-shaped nanograins are present at the bottom of this section (adjacent to the submicro-sized section). Figs. 11 and 12 show typical cross-sectional observations of the micro￾structure in this section. In Fig. 11(a), lamellar grains of 10–50 nm thick and 50–100 nm long are observed. It is worth noting that inside the inner of these lamellar nanocrystallites, there exist even smaller equiaxed nano-grains. In Fig. 11(b), one can see roughly equiaxed nanocrystallites with sizes of about 30 nm. The SAED pattern indicates small misorientations among these nanocrystal￾lites. Evidently, these equiaxed nanograins were

N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 4611 200nm 200nm Fig.10.A cross-sectional TEM image in the submicro-sized section(28 um deep)showing equiaxed grains separated by Fig.9.A cross-sectional TEM image in the submicro-sized large misorientaion boundaries. section showing equiaxed subgrains with different dislocation configurations and boundaries(solid triangles for DDWs and boundaries along the short axis.The random orien- DTs,open triangles for subboundaries). tations of the nanocrystallites may be resulted from formation of high-angle grain boundary and/or grain rotation under large strains. formed by breaking up the nano-sized lamellar A close observation of the nanostructure in the grains (or cells)due to dislocation accumulation top surface layer was performed by using HRTEM, and subboundary development,as in the case of as depicted in Fig.13.The nano-sized crystallites submicro-sized section.Misorientations between with different misorientations were clearly ident- the neighboring nanocrystallites gradually increase ified,including large angles(nanocrystallites a and to large angles with further straining.Then,equi- b)and small ones (nanocrystallites c and d,about axed nanocrystallites with random orientations can 5).From the plane-view (Fig.3 and Fig.13)and be achieved.Fig.12 shows a cross-sectional TEM the cross-sectional view (Fig.12)of the top surface observation at the very top surface layer.One may layer,one may conclude that equixed nano-sized see equiaxed nanocrystallites of about 10-20 nm grains are formed in the top surface layer in the in size,and their crystallographic orientations are Fe sample. random,as indicated by the SAED pattern.From the dark-field TEM image in Fig.12,it is interest- ing to find that the nanocrystallites seem to be 4.Discussion arranged in parallel arrays.The nanocrystallite arrays seem to be formed by cutting nano-sized Microstructural investigations revealed that lamellar cells by means of development of more ultrafine-grained structures (from nano-to micro-

N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 4611 Fig. 9. A cross-sectional TEM image in the submicro-sized section showing equiaxed subgrains with different dislocation configurations and boundaries (solid triangles for DDWs and DTs, open triangles for subboundaries). formed by breaking up the nano-sized lamellar grains (or cells) due to dislocation accumulation and subboundary development, as in the case of submicro-sized section. Misorientations between the neighboring nanocrystallites gradually increase to large angles with further straining. Then, equi￾axed nanocrystallites with random orientations can be achieved. Fig. 12 shows a cross-sectional TEM observation at the very top surface layer. One may see equiaxed nanocrystallites of about 10–20 nm in size, and their crystallographic orientations are random, as indicated by the SAED pattern. From the dark-field TEM image in Fig. 12, it is interest￾ing to find that the nanocrystallites seem to be arranged in parallel arrays. The nanocrystallite arrays seem to be formed by cutting nano-sized lamellar cells by means of development of more Fig. 10. A cross-sectional TEM image in the submicro-sized section (28 µm deep) showing equiaxed grains separated by large misorientaion boundaries. boundaries along the short axis. The random orien￾tations of the nanocrystallites may be resulted from formation of high-angle grain boundary and/or grain rotation under large strains. A close observation of the nanostructure in the top surface layer was performed by using HRTEM, as depicted in Fig. 13. The nano-sized crystallites with different misorientations were clearly ident￾ified, including large angles (nanocrystallites a and b) and small ones (nanocrystallites c and d, about 5°). From the plane-view (Fig. 3 and Fig. 13) and the cross-sectional view (Fig. 12) of the top surface layer, one may conclude that equixed nano-sized grains are formed in the top surface layer in the Fe sample. 4. Discussion Microstructural investigations revealed that ultrafine-grained structures (from nano- to micro-

4612 N.R.Tao et al.Acta Materialia 50 (2002)4603-4616 (a) 100nm b Fig.13.A plane-view HRTEM image of the nanostructure in the top surface layer. sized)were formed in the surface layer of the Fe sample during the SMA treatment.Based on the microstructure features observed in various sec- tions with different strains in the deformed surface layer,one may find that the following elemental 100nm processes are involved in the grain refinement pro- cess: Fig.11.Cross-sectional TEM images in the top surface layer showing:(a)lamellar nanocrystallites;(b)equiaxed nanocrys- 1.development of DDWs and DTs in original tallites with small angle misorientations. grains and in the refined cells (under further straining)as well; 2.transformation of DDWs and DTs into subboun- daries with small misorientations separating individual cells or subgrains; 3.evolution of subboundaries to highly misori- ented grain boundaries. The grain refinement mechanism can be sche- matically illustrated in Fig.14,in which each pro- cess will be discussed in terms of the experi- mental observations. 4.1.Development of DDWs and DTs 100nm 100nm In order to accommodate plastic strains in Fig.12.(a)A bright-field and (b)a dark-field cross-sectional polycrystalline materials,various dislocation TEM images showing nanocrystallites in the very top surface activities are normally motivated,including slid- layer. ing,accumulation,interaction,tangling,and spatial rearrangement.In the SMA treated Fe sample,dis- location activities lead to formation of DDWs and DTs in original grains of the surface layer.Devel-

4612 N.R. Tao et al. / Acta Materialia 50 (2002) 4603–4616 Fig. 11. Cross-sectional TEM images in the top surface layer showing: (a) lamellar nanocrystallites; (b) equiaxed nanocrys￾tallites with small angle misorientations. Fig. 12. (a) A bright-field and (b) a dark-field cross-sectional TEM images showing nanocrystallites in the very top surface layer. Fig. 13. A plane-view HRTEM image of the nanostructure in the top surface layer. sized) were formed in the surface layer of the Fe sample during the SMA treatment. Based on the microstructure features observed in various sec￾tions with different strains in the deformed surface layer, one may find that the following elemental processes are involved in the grain refinement pro￾cess: 1. development of DDWs and DTs in original grains and in the refined cells (under further straining) as well; 2. transformation of DDWs and DTs into subboun￾daries with small misorientations separating individual cells or subgrains; 3. evolution of subboundaries to highly misori￾ented grain boundaries. The grain refinement mechanism can be sche￾matically illustrated in Fig. 14, in which each pro￾cess will be discussed in terms of the experi￾mental observations. 4.1. Development of DDWs and DTs In order to accommodate plastic strains in polycrystalline materials, various dislocation activities are normally motivated, including slid￾ing, accumulation, interaction, tangling, and spatial rearrangement. In the SMA treated Fe sample, dis￾location activities lead to formation of DDWs and DTs in original grains of the surface layer. Devel-