Act MATERIALIA Pergamon Acta mater.49(2001)1497-1505 www.elsevier.com/locate/actamat MICROSTRUCTURES AND DISLOCATION CONFIGURATIONS IN NANOSTRUCTURED Cu PROCESSED BY REPETITIVE CORRUGATION AND STRAIGHTENING J.Y.HUANGi,Y.T.ZHU,H.JIANG and T.C.LOWE Materials Science and Technology Division,MS G 755,Los Alamos National Laboratory,Los Alamos, NM 87545.USA Received 6 November 2000:accepted 5 February 2001 Abstract-The microstructures and dislocation configurations in nanostructured Cu processed by a new tech- nique,repetitive corrugation and straightening (RCS),were studied using transmission electron microcopy (TEM)and high resolution TEM.Most dislocations belong to 60 type and tend to pile up along the (111] slip planes.Microstructural features including low-angle grain boundaries(GBs),high-angle GBs,and equilib- rium and non-equilibrium GBs and subgrain boundaries were observed.Dislocation structures at an intermedi- ate deformation strain were studied to investigate the microstructural evolutions,which revealed some unique microstructural features such as isolated dislocation cell (IDC),dislocation tangle zones (DTZs),and uncon- densed dislocation walls (UDWs).2001 Acta Materialia Inc.Published by Elsevier Science Ltd.All rights reserved. Keywords:Repetitive corrugation straightening;Microstructure;Transmission electron microscopy (TEM): Copper 1.INTRODUCTION has been paid to alternative procedures of introducing ultrafine grains in materials by severe plastic defor- Many methods have been used to synthesize materials mation (SPD)[9-121. with ultrafine grain sizes (10-1000 nm),including inert gas condensation [1],high-energy ball milling One of the SPD variants,equal-channel angular pressing (ECAP),has been used to refine bulk, [2],sliding wear [3],etc.These techniques are attract- ive for producing powders with grain sizes below coarse-grained metals and alloys to grain sizes rang- 100 nm,but cannot be used to make bulk samples. ing from <0.I to I um [9-12].However,ECAP is To consolidate the nanometer-sized powders into bulk difficult to scale up to process volumes of materials materials,high pressure and moderate temperature are much larger than the 20x20x100 mm3 samples that usually needed.Grains might grow during consoli- are typically produced today.Furthermore,current dation,making the bulk materials partially or com- implementations of ECAP are discontinuous,requir- pletely lose the nanocharacteristics.It is usually ing labor intensive handling of the work-piece impossible to completely eliminate porosity,even in between process steps.These difficulties in fabricat- materials consolidated under very high pressure and ing bulk,nanostructured materials have been substan- temperature.In addition,nanopowders are very sus- tial road-blocks to the structural applications of nano- ceptible to oxidation and absorb large quantities of structured materials.Other SPD techniques that have impurities such as O2.H2 and N2,making it difficult been reported in the literature include multipass-coin- to obtain clean bulk materials.The porosity as well forge (MCF)[13]and multi-axis deformation [14]. as impurities significantly affect the mechanical Both of them have certain advantages over the ECAP properties of the bulk materials,often making them process.However,they also employ batch processing, brittle [4-8].These problems prevent us from study- which is not efficient for large-scale production. ing the intrinsic properties of bulk nanomaterials.As Recently,we have developed a new technique, a consequence of these difficulties,much attention repetitive corrugation and straightening (RCS),that can not only create bulk nanostructured materials free of contamination and porosity,but can also be easily To whom all correspondence should be addressed.Tel.: adapted to large-scale industrial production [15].In +1-505-665-0835:fax:+1-505-667-2264. the RCS process,a work-piece is repetitively bent and E-mail address:jyhuang@lanl.gov (J.Huang) straightened without significantly changing the cross- 1359-6454/01/$20.00 2001 Acta Materialia Inc.Published by Elsevier Science Ltd.All rights reserved. P:S1359-6454(01)00069-6
Acta mater. 49 (2001) 1497–1505 www.elsevier.com/locate/actamat MICROSTRUCTURES AND DISLOCATION CONFIGURATIONS IN NANOSTRUCTURED Cu PROCESSED BY REPETITIVE CORRUGATION AND STRAIGHTENING J. Y. HUANG†, Y. T. ZHU, H. JIANG and T. C. LOWE Materials Science and Technology Division, MS G 755, Los Alamos National Laboratory, Los Alamos, NM 87545, USA ( Received 6 November 2000; accepted 5 February 2001 ) Abstract—The microstructures and dislocation configurations in nanostructured Cu processed by a new technique, repetitive corrugation and straightening (RCS), were studied using transmission electron microcopy (TEM) and high resolution TEM. Most dislocations belong to 60° type and tend to pile up along the {111} slip planes. Microstructural features including low-angle grain boundaries (GBs), high-angle GBs, and equilibrium and non-equilibrium GBs and subgrain boundaries were observed. Dislocation structures at an intermediate deformation strain were studied to investigate the microstructural evolutions, which revealed some unique microstructural features such as isolated dislocation cell (IDC), dislocation tangle zones (DTZs), and uncondensed dislocation walls (UDWs). 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Repetitive corrugation straightening; Microstructure; Transmission electron microscopy (TEM); Copper 1. INTRODUCTION Many methods have been used to synthesize materials with ultrafine grain sizes (10–1000 nm), including inert gas condensation [1], high-energy ball milling [2], sliding wear [3], etc. These techniques are attractive for producing powders with grain sizes below 100 nm, but cannot be used to make bulk samples. To consolidate the nanometer-sized powders into bulk materials, high pressure and moderate temperature are usually needed. Grains might grow during consolidation, making the bulk materials partially or completely lose the nanocharacteristics. It is usually impossible to completely eliminate porosity, even in materials consolidated under very high pressure and temperature. In addition, nanopowders are very susceptible to oxidation and absorb large quantities of impurities such as O2, H2 and N2, making it difficult to obtain clean bulk materials. The porosity as well as impurities significantly affect the mechanical properties of the bulk materials, often making them brittle [4–8]. These problems prevent us from studying the intrinsic properties of bulk nanomaterials. As a consequence of these difficulties, much attention † To whom all correspondence should be addressed. Tel.: +1-505-665-0835; fax: +1-505-667-2264. E-mail address: jyhuang@lanl.gov (J. Huang) 1359-6454/01/$20.00 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S13 59-6454(01)00069-6 has been paid to alternative procedures of introducing ultrafine grains in materials by severe plastic deformation (SPD) [9–12]. One of the SPD variants, equal-channel angular pressing (ECAP), has been used to refine bulk, coarse-grained metals and alloys to grain sizes ranging from 0.1 to 1 µm [9–12]. However, ECAP is difficult to scale up to process volumes of materials much larger than the 20×20×100 mm3 samples that are typically produced today. Furthermore, current implementations of ECAP are discontinuous, requiring labor intensive handling of the work-piece between process steps. These difficulties in fabricating bulk, nanostructured materials have been substantial road-blocks to the structural applications of nanostructured materials. Other SPD techniques that have been reported in the literature include multipass-coinforge (MCF) [13] and multi-axis deformation [14]. Both of them have certain advantages over the ECAP process. However, they also employ batch processing, which is not efficient for large-scale production. Recently, we have developed a new technique, repetitive corrugation and straightening (RCS), that can not only create bulk nanostructured materials free of contamination and porosity, but can also be easily adapted to large-scale industrial production [15]. In the RCS process, a work-piece is repetitively bent and straightened without significantly changing the cross-
1498 HUANG et al:NANOSTRUCTURED Cu section geometry of the work-piece,during which grain size is desired to effectively demonstrate the large plastic strains are imparted into the materials, grain-refinement capability of the RCS process.A which leads to the refinement of the microstructure. basic RCS cycle consists of two steps:corrugation Hansen and coworkers [16-27]have systematically and straightening.The corrugation is carried out in a studied the evolution of microstructures and defined die set as shown in Fig.2(a),which is the discontinu- microstructural features in rolled face-centered cubic ous version of the RCS process.It was used to simu- (fcc)metals with medium to high stacking fault ener- late the continuous version of the RCS process as gies,such as Al and Cu.However,the deformation shown in Fig.2(b).It is obvious that the continuous mode in the RCS process is different from that of RCS process can be easily adapted to a rolling mill rolling and is expected to result in different micro- for industrial production of nanostructured metals and structural evolution and consequently different micro- alloys.The straightening is accomplished by pressing structures. the corrugated work-piece between two flat platens. A controversial microstructural feature in nanos- It is well known that lower deformation temperature tructured materials processed by SPD techniques is impedes dynamic recovery and consequently non-equilibrium grain boundary (non-equilibrium improves the grain-refinement efficiency [35].There- GB)[28,29].Valiev et al.[28]defined it as GB that fore,the copper bar was immersed in liquid nitrogen contains extrinsic dislocations that are not needed to for 3 min before each RCS cycle.It took about 1 min accommodate the misorientation across the GB.The to take the sample out of liquid nitrogen,place it in extrinsic dislocations are usually lattice dislocations the die and start pressing it.The sample temperature trapped at the GB.They cause lattice distortion near was unknown but is expected to rise with time during the GB and increase the GB energy [30].Although the RCS process.This corrugation-straightening cycle the non-equilibrium GB has been mentioned by many was repeated 14 times with 90 rotations along the researchers [12,30-34],it has not been directly longitudinal axis of the sample between consecutive proved experimentally and has been controversial. cycles.Lubricant was used to reduce friction between The objectives of this work are:(1)to study the the work-piece and the die,although some amount of microstructural features and dislocation configur- stretching was still present.The microstructure of the ations in nanostructured Cu processed by RCS:(2)to investigate the microstructural evolutions and grain- refinement mechanisms through TEM observations of (a) Press dislocation structures at an intermediate deformation strain;and (3)to clarify the existence of non-equilib- rium GBs produced by SPD. 2.EXPERIMENTAL PROCEDURES A high purity (99.99 at.%)copper bar with 6x6x50 mm3 in dimension was used in this study.It was annealed at 900C for I h to increase its average Work piece grain size to about 765 um (see Fig.1).The large (b) 500m Fig.2.(a)Die set up for discontinuous RCS process and (b) Fig.1.As-annealed copper has an average grain size of set up for continuous RCS process that can be easily adapted 765μm. to large-scale industrial production
1498 HUANG et al.: NANOSTRUCTURED Cu section geometry of the work-piece, during which large plastic strains are imparted into the materials, which leads to the refinement of the microstructure. Hansen and coworkers [16–27] have systematically studied the evolution of microstructures and defined microstructural features in rolled face-centered cubic (fcc) metals with medium to high stacking fault energies, such as Al and Cu. However, the deformation mode in the RCS process is different from that of rolling and is expected to result in different microstructural evolution and consequently different microstructures. A controversial microstructural feature in nanostructured materials processed by SPD techniques is non-equilibrium grain boundary (non-equilibrium GB) [28, 29]. Valiev et al. [28] defined it as GB that contains extrinsic dislocations that are not needed to accommodate the misorientation across the GB. The extrinsic dislocations are usually lattice dislocations trapped at the GB. They cause lattice distortion near the GB and increase the GB energy [30]. Although the non-equilibrium GB has been mentioned by many researchers [12, 30–34], it has not been directly proved experimentally and has been controversial. The objectives of this work are: (1) to study the microstructural features and dislocation configurations in nanostructured Cu processed by RCS; (2) to investigate the microstructural evolutions and grainrefinement mechanisms through TEM observations of dislocation structures at an intermediate deformation strain; and (3) to clarify the existence of non-equilibrium GBs produced by SPD. 2. EXPERIMENTAL PROCEDURES A high purity (99.99 at.%) copper bar with 6×6×50 mm3 in dimension was used in this study. It was annealed at 900°C for 1 h to increase its average grain size to about 765 µm (see Fig. 1). The large Fig. 1. As-annealed copper has an average grain size of 765 µm. grain size is desired to effectively demonstrate the grain-refinement capability of the RCS process. A basic RCS cycle consists of two steps: corrugation and straightening. The corrugation is carried out in a die set as shown in Fig. 2(a), which is the discontinuous version of the RCS process. It was used to simulate the continuous version of the RCS process as shown in Fig. 2(b). It is obvious that the continuous RCS process can be easily adapted to a rolling mill for industrial production of nanostructured metals and alloys. The straightening is accomplished by pressing the corrugated work-piece between two flat platens. It is well known that lower deformation temperature impedes dynamic recovery and consequently improves the grain-refinement efficiency [35]. Therefore, the copper bar was immersed in liquid nitrogen for 3 min before each RCS cycle. It took about 1 min to take the sample out of liquid nitrogen, place it in the die and start pressing it. The sample temperature was unknown but is expected to rise with time during the RCS process. This corrugation-straightening cycle was repeated 14 times with 90° rotations along the longitudinal axis of the sample between consecutive cycles. Lubricant was used to reduce friction between the work-piece and the die, although some amount of stretching was still present. The microstructure of the Fig. 2. (a) Die set up for discontinuous RCS process and (b) set up for continuous RCS process that can be easily adapted to large-scale industrial production
HUANG et al:NANOSTRUCTURED Cu 1499 deformed sample was characterized by transmission electron microscopy (TEM)as well as high-resolution TEM (HRTEM).The HRTEM was carried out in a JEOL 3000 FEG electron microscope operated at 300 kV.The point-to-point resolution is about 1.8 A. TEM and HRTEM samples were prepared by jet elec- tro-polishing at room temperature.The electrolyte consists of 33%orthophosphoric acid and 67%water. To enhance the image contrast,most of the HREM images were reconstructed from Fast Fourier Trans- formation,during which the diffuse scattering from the background or inelastic scattering was filtered. DTZ 3.EXPERIMENTAL RESULTS 3.1.Microstructures and dislocations in nanostruc- tured Cu processed for 14 RCS passes Figure 3(a)is a TEM micrograph showing that individual grains were produced with sizes ranging from less than 100 nm to a few hundred nanometers, 100nm separated by high-angle grain boundaries(high-angle GBs).Most grains are heavily strained and contain high density of dislocations.The corresponding elec- tron diffraction pattern (EDP)in Fig.3(b)exhibits diffraction rings,indicating a polycrystalline struc- ture.The diffraction rings show significant 011 tex- ture.Figure 4(a)shows a TEM micrograph of a grain with a diameter of about 500 nm.A number of fine structures were observed in the interior of the grain. As pointed out by two arrowheads,an array of dislo- 3nm cations piled up along the (111)plane.Consequently, two subgrains (denoted by I and 2)with a misorien- tation of about 1 were produced(measured from the Fig.4.(a)A TEM micrograph showing fine deformation struc- tures in a grain.The numbers 1-3 denote three subgrains:the HRTEM image,not shown here).The dislocations are two arrowheads point out an array of dislocations:the four stars mostly 60 type (as shown later in Fig.4(c))and are mark a low-angle GB:the white circle marks a dislocation tan- glissible along the (111}planes. gle zone (DTZ):the white square marks a transition from DTZ A low-angle GB was also found in this grain,as to dislocation cells.(b)An HRTEM image of the low-angle GB pointed out by the four stars in (a).(c)A Fourier filtered marked by four stars in Fig.4(a).The low-angle GB HRTEM image of a 60 dislocation.A Burgers circuit was was formed by the accumulation of a number of gliss- drawn to enclose the dislocation core marked by a "T".The ile dislocations.It is not edge-on but overlapped,as electron beam and the dislocation line is parallel to [11O].and revealed by the periodic Moire Fringes.Figure 4(b) the Burgers vector b =1/2[011]or 1/2[101]. is an HRTEM image of a local region of this low- angle GB.The misorientation of the two grains is about 5.The spacing of the Moire pattern can be d=2.08 A.D is calculated to be 23.84 A.which calculated using the formula:D=d/o,where d is the agrees well with the experimental value of 24 A,as measured from Fig.4(b). lattice spacing and a is the rotation angle.For Cu. Dislocation cell structure was also observed in subgrain 3 in Fig.4(a).These cells may form individ- ual subgrains upon further plastic straining.Dislo- cation tangling was frequently observed in the interior of grains,as marked by a white circle in Fig.4(a). where the grain is heavily strained.We shall refer such a region as dislocation-tangle zone (DTZ).Fig- ure 4(c)shows a Fourier filtered HRTEM image of a 60 dislocation which was frequently observed in RCS-deformed Cu.Assuming the electron beam and the dislocation line is parallel to [110].the Burgers vector of the dislocation is determined to be 1/2[101] Fig.3.TEM micrographs showing:(a)nanostructured Cu pro- or 1/210111.which has an angle of 60(or 120)with duced by the RCS process:and (b)the corresponding EDP. respect to the dislocation line.For this reason,the
HUANG et al.: NANOSTRUCTURED Cu 1499 deformed sample was characterized by transmission electron microscopy (TEM) as well as high-resolution TEM (HRTEM). The HRTEM was carried out in a JEOL 3000 FEG electron microscope operated at 300 kV. The point-to-point resolution is about 1.8 A˚ . TEM and HRTEM samples were prepared by jet electro-polishing at room temperature. The electrolyte consists of 33% orthophosphoric acid and 67% water. To enhance the image contrast, most of the HREM images were reconstructed from Fast Fourier Transformation, during which the diffuse scattering from the background or inelastic scattering was filtered. 3. EXPERIMENTAL RESULTS 3.1. Microstructures and dislocations in nanostructured Cu processed for 14 RCS passes Figure 3(a) is a TEM micrograph showing that individual grains were produced with sizes ranging from less than 100 nm to a few hundred nanometers, separated by high-angle grain boundaries (high-angle GBs). Most grains are heavily strained and contain high density of dislocations. The corresponding electron diffraction pattern (EDP) in Fig. 3(b) exhibits diffraction rings, indicating a polycrystalline structure. The diffraction rings show significant 011 texture. Figure 4(a) shows a TEM micrograph of a grain with a diameter of about 500 nm. A number of fine structures were observed in the interior of the grain. As pointed out by two arrowheads, an array of dislocations piled up along the (111) plane. Consequently, two subgrains (denoted by 1 and 2) with a misorientation of about 1° were produced (measured from the HRTEM image, not shown here). The dislocations are mostly 60° type (as shown later in Fig. 4(c)) and are glissible along the {111} planes. A low-angle GB was also found in this grain, as marked by four stars in Fig. 4(a). The low-angle GB was formed by the accumulation of a number of glissile dislocations. It is not edge-on but overlapped, as revealed by the periodic Moire´ Fringes. Figure 4(b) is an HRTEM image of a local region of this lowangle GB. The misorientation of the two grains is about 5°. The spacing of the Moire´ pattern can be calculated using the formula: D = d/a, where d is the lattice spacing and a is the rotation angle. For Cu, Fig. 3. TEM micrographs showing: (a) nanostructured Cu produced by the RCS process; and (b) the corresponding EDP. Fig. 4. (a) A TEM micrograph showing fine deformation structures in a grain. The numbers 1–3 denote three subgrains; the two arrowheads point out an array of dislocations; the four stars mark a low-angle GB; the white circle marks a dislocation tangle zone (DTZ); the white square marks a transition from DTZ to dislocation cells. (b) An HRTEM image of the low-angle GB pointed out by the four stars in (a). (c) A Fourier filtered HRTEM image of a 60° dislocation. A Burgers circuit was drawn to enclose the dislocation core marked by a “T”. The electron beam and the dislocation line is parallel to [11¯0], and the Burgers vector b = 1/2[011] or 1/2[101]. d(111) = 2.08 A˚ . D is calculated to be 23.84 A˚ , which agrees well with the experimental value of 24 A˚ , as measured from Fig. 4(b). Dislocation cell structure was also observed in subgrain 3 in Fig. 4(a). These cells may form individual subgrains upon further plastic straining. Dislocation tangling was frequently observed in the interior of grains, as marked by a white circle in Fig. 4(a), where the grain is heavily strained. We shall refer such a region as dislocation-tangle zone (DTZ). Figure 4(c) shows a Fourier filtered HRTEM image of a 60° dislocation which was frequently observed in RCS-deformed Cu. Assuming the electron beam and the dislocation line is parallel to [110], the Burgers vector of the dislocation is determined to be 1/2[101] or 1/2[011], which has an angle of 60° (or 120°) with respect to the dislocation line. For this reason, the
1500 HUANG et al:NANOSTRUCTURED Cu dislocation is referred as a 60 dislocation.In many subgrain boundary as non-equilibrium subgrain cases,such a dislocation is dissociated or extended to boundary. more than 10 atomic planes. Figure 6(a)shows a Fourier filtered HRTEM image Figure 5(a)shows another example of subgrain of a low-angle GB,which is delineated by periodic generation from larger grains.The parallelogram- dislocations and Moire Fringes.The GB plane is shaped subgrain with a size of about 250 nm is delin- curved and changes from the (55 12)plane to the eated clearly by the dense-dislocation walls (DDWs) (002)plane.The corresponding EDP (Fig.6(b)) [18,19],which are almost parallel to the two sets of shows that the two grains are misoriented for about (111)planes.This subgrain is clearly inside a larger 9.HRTEM images from the upper-left and lower- grain and is isolated from other subgrains.We shall right part of the low-angle GB(see the framed areas) call it isolated subgrains since its boundary does not are shown in Fig.6(c)and (e),respectively.Figure meet with other subgrain boundaries.The dislocation 6(d)is a structural model corresponding to the low- density is higher inside the subgrain than outside it.angle GB in Fig.6(c).From this model,it is seen that Some of the dislocations are forming cell structures two types of dislocations with Burgers vectors inside the subgrain.Figure 5(b)is a Fourier filtered b=1/2[1011 and b2 1/2[1011.hereafter referred as HRTEM image of the dislocation wall at the point type 1 and type 2 dislocations,respectively,are marked by an arrowhead in Fig.5(a).The dislocation needed to accommodate the geometrical misorien- density is estimated as 3x1017 m-2 at the DDW tation.In other words,these dislocations are geo- subgrain boundary.The estimation is based on the metrically necessary.Valiev et al.[28]referred these Fourier filtered HREM image shown in Fig.5(b).The geometrically necessary dislocations as intrinsic dis- number of dislocations was counted,and divided by locations.According to Fig.6(d),the spacing of type the area in this figure.Such an estimation assumes I dislocations is about 18 A,which is consistent with that the dislocation line goes straight from the top of that measured from Fig.6(c).However,there are the grain through the bottom of it,which may not be three more type 2 dislocations in Fig.6(c)than in(d), the case for most dislocations,since most of them are which indicates that three extrinsic (or non-geometri- curved,or they may terminate in the grain interior. cally necessary)dislocations exist at the GB shown Therefore such estimation is only qualitative but not in Fig.6(c).Therefore,this segment of low-angle GB quantitative.Interstitial loops (marked by black is in a high energy configuration and should be called circles)and vacancy loops (marked by white circles) non-equilibrium grain boundary also exist.The dislocations are again mostly 60 type In Fig.6(e),dislocations are periodically spaced. ones.In addition,the lattice planes near the cell walls The Burgers vector is determined as 1/2[101].The are heavily distorted.The width of the subgrain dislocation spacing in a low-angle GB can be calcu- boundary is about 10 nm.The misorientation across lated using the formula:D b/e,where b is the Burg- the subgrain boundary is measured as about 5.There ers vector of the GB dislocation and 6 the rotation are significantly more dislocations than required to angle of the two grains.The calculated dislocation accommodate the misorientation.These dislocations spacing is 20 A,which is in reasonable agreement are not arranged in the lowest-energy dislocation with experimentally measured values of 22 A in Fig. structure (LEDS)[17,19,36],which makes the 6(e).No extrinsic dislocation is found in Fig.6(e). subgrain boundary unstable.We shall refer such This segment GB is equilibrium GB. Besides the 60 dislocations,other dislocations such as screw dislocations and Frank dislocations were also frequently observed.Figure 7(a)shows a number of dislocations in a grain.The Fourier filtered HRTEM image shown in Fig 7(b)reveals that these dislocations are Frank dislocations.The Burgers vec- tor was determined as 1/2[110].which is in (110) plane.Therefore,they are immobile or sessile dislo- cations. 20m 3.2.Deformation microstructures at intermediate plastic strain Fig.5.(a)A TEM micrograph of a subgrain;inset is an To investigate the microstructural evolutions and HRTEM image from the subgrain showing the crystalline planes.Note that the subgrain is delineated by DDWs which grain-refinement mechanisms,deformation structures are almost parallel to two sets of (111)planes.(b)A Fourier at an intermediate deformation strain (6 RCS passes) filtered HRTEM image from the DDW as pointed out by an were studied.Shown in Fig.8 is a TEM micrograph arrowhead in (a).The electron beam and the dislocation line that depicts many microstructural features.The is parallel to [1fo],and the Burgers vector b 1/2[011]or description of deformation microstructures is contro- 1/2[101].The dislocations are all 60 type.The longer arrow points out the grain boundary orientation.The black circles versial and is confusing in the literature [37].In a mark interstitial loops and the white circles mark vacancy series of recent papers [16-27].Hansen and cowork- loops. ers systematically studied the evolution of microstruc-
1500 HUANG et al.: NANOSTRUCTURED Cu dislocation is referred as a 60° dislocation. In many cases, such a dislocation is dissociated or extended to more than 10 atomic planes. Figure 5(a) shows another example of subgrain generation from larger grains. The parallelogramshaped subgrain with a size of about 250 nm is delineated clearly by the dense-dislocation walls (DDWs) [18,19], which are almost parallel to the two sets of {111} planes. This subgrain is clearly inside a larger grain and is isolated from other subgrains. We shall call it isolated subgrains since its boundary does not meet with other subgrain boundaries. The dislocation density is higher inside the subgrain than outside it. Some of the dislocations are forming cell structures inside the subgrain. Figure 5(b) is a Fourier filtered HRTEM image of the dislocation wall at the point marked by an arrowhead in Fig. 5(a). The dislocation density is estimated as 3×1017 m2 at the DDW subgrain boundary. The estimation is based on the Fourier filtered HREM image shown in Fig. 5(b). The number of dislocations was counted, and divided by the area in this figure. Such an estimation assumes that the dislocation line goes straight from the top of the grain through the bottom of it, which may not be the case for most dislocations, since most of them are curved, or they may terminate in the grain interior. Therefore such estimation is only qualitative but not quantitative. Interstitial loops (marked by black circles) and vacancy loops (marked by white circles) also exist. The dislocations are again mostly 60° type ones. In addition, the lattice planes near the cell walls are heavily distorted. The width of the subgrain boundary is about 10 nm. The misorientation across the subgrain boundary is measured as about 5°. There are significantly more dislocations than required to accommodate the misorientation. These dislocations are not arranged in the lowest-energy dislocation structure (LEDS) [17, 19, 36], which makes the subgrain boundary unstable. We shall refer such Fig. 5. (a) A TEM micrograph of a subgrain; inset is an HRTEM image from the subgrain showing the crystalline planes. Note that the subgrain is delineated by DDWs which are almost parallel to two sets of {111} planes. (b) A Fourier filtered HRTEM image from the DDW as pointed out by an arrowhead in (a). The electron beam and the dislocation line is parallel to [11¯0], and the Burgers vector b = 1/2[011] or 1/2[101]. The dislocations are all 60° type. The longer arrow points out the grain boundary orientation. The black circles mark interstitial loops and the white circles mark vacancy loops. subgrain boundary as non-equilibrium subgrain boundary. Figure 6(a) shows a Fourier filtered HRTEM image of a low-angle GB, which is delineated by periodic dislocations and Moire´ Fringes. The GB plane is curved and changes from the (5 5 12) plane to the (002) plane. The corresponding EDP (Fig. 6(b)) shows that the two grains are misoriented for about 9°. HRTEM images from the upper-left and lowerright part of the low-angle GB (see the framed areas) are shown in Fig. 6(c) and (e), respectively. Figure 6(d) is a structural model corresponding to the lowangle GB in Fig. 6(c). From this model, it is seen that two types of dislocations with Burgers vectors b1 = 1/2[101] and b2 = 1/2[101], hereafter referred as type 1 and type 2 dislocations, respectively, are needed to accommodate the geometrical misorientation. In other words, these dislocations are geometrically necessary. Valiev et al. [28] referred these geometrically necessary dislocations as intrinsic dislocations. According to Fig. 6(d), the spacing of type 1 dislocations is about 18 A˚ , which is consistent with that measured from Fig. 6(c). However, there are three more type 2 dislocations in Fig. 6(c) than in (d), which indicates that three extrinsic (or non-geometrically necessary) dislocations exist at the GB shown in Fig. 6(c). Therefore, this segment of low-angle GB is in a high energy configuration and should be called non-equilibrium grain boundary. In Fig. 6(e), dislocations are periodically spaced. The Burgers vector is determined as 1/2[101]. The dislocation spacing in a low-angle GB can be calculated using the formula: D = b/q, where b is the Burgers vector of the GB dislocation and q the rotation angle of the two grains. The calculated dislocation spacing is 20 A˚ , which is in reasonable agreement with experimentally measured values of 22 A˚ in Fig. 6(e). No extrinsic dislocation is found in Fig. 6(e). This segment GB is equilibrium GB. Besides the 60° dislocations, other dislocations such as screw dislocations and Frank dislocations were also frequently observed. Figure 7(a) shows a number of dislocations in a grain. The Fourier filtered HRTEM image shown in Fig 7(b) reveals that these dislocations are Frank dislocations. The Burgers vector was determined as 1/2[1¯10], which is in (11¯0) plane. Therefore, they are immobile or sessile dislocations. 3.2. Deformation microstructures at intermediate plastic strain To investigate the microstructural evolutions and grain-refinement mechanisms, deformation structures at an intermediate deformation strain (6 RCS passes) were studied. Shown in Fig. 8 is a TEM micrograph that depicts many microstructural features. The description of deformation microstructures is controversial and is confusing in the literature [37]. In a series of recent papers [16–27], Hansen and coworkers systematically studied the evolution of microstruc-
HUANG et al:NANOSTRUCTURED Cu 1501 55 002 nm 2nm Fig.6.(a)A TEM micrograph of a low-angle GB.The GB plane is curved from(55 12)to (002)plane.(b) An EDP corresponding to (a).The grain misorientation measured from the EDP is 9.(c)and (d)A Fourier filtered HRTEM image and a structural model of the upper-left of the low-angle GB shown in (a)(marked by a black frame).Burgers vectors are b =1/2[101]and b2=1/2[101].(e)A Fourier filtered HRTEM image from the lower-right part of the low-angle GB shown in (a).The Burgers vectors b of the dislocations is 1/2[101].The unit cells of the images are also outlined. As shown.the deformation structure consists of 110 dislocation cells and cell-blocks (CBs)[20].Within each CB,the deformation is accommodated by one set of slip systems that are different from those in neighboring CBs.The boundary between the neighb- oring CBs is called geometrically necessary boundary (GNB)since the boundary is necessary to accommo- date the glide-induced lattice rotation in the adjoining a CBs [20].Figure 8 shows that GNBs may be parallel to the (111}slip planes or other crystallographic Fig.7.(a)A TEM micrograph showing the high density of planes.The dislocation cell boundaries are called dislocations in a grain.(b)A Fourier filtered HRTEM image showing the dislocation type in (a).The Burgers vector determ- incidental boundaries since they are generated by the ined from the Burgers circuit is b 1/2[110]. statistical mutual trappings of glide dislocations. Hansen and coworkers [16,17,19]observed two types of features on the GNBs in the rolling-induced tures and defined microstructural features in rolled fcc microstructure.One is DDWs.The other is metals with medium to high stacking fault energies, microbands(MBs)that are composed of small pan- such as Al and Cu.In Fig.8,we try to follow their cake-shaped cells(SPCs).The definition of MBs has nomenclatures in defining the microstructural fea- been controversial and confusing since the nomencla- tures,but find it necessary to make our own ture has been used to describe the long,thin plate-like nomenclatures in several occasions. regions observed in deformed copper [37-39].The
HUANG et al.: NANOSTRUCTURED Cu 1501 Fig. 6. (a) A TEM micrograph of a low-angle GB. The GB plane is curved from (5 5 12) to (002) plane. (b) An EDP corresponding to (a). The grain misorientation measured from the EDP is 9°. (c) and (d) A Fourier filtered HRTEM image and a structural model of the upper-left of the low-angle GB shown in (a) (marked by a black frame). Burgers vectors are b1 = 1/2[101] and b2 = 1/2[101¯]. (e) A Fourier filtered HRTEM image from the lower-right part of the low-angle GB shown in (a). The Burgers vectors b of the dislocations is 1/2[101]. The unit cells of the images are also outlined. Fig. 7. (a) A TEM micrograph showing the high density of dislocations in a grain. (b) A Fourier filtered HRTEM image showing the dislocation type in (a). The Burgers vector determined from the Burgers circuit is b = 1/2[1¯10]. tures and defined microstructural features in rolled fcc metals with medium to high stacking fault energies, such as Al and Cu. In Fig. 8, we try to follow their nomenclatures in defining the microstructural features, but find it necessary to make our own nomenclatures in several occasions. As shown, the deformation structure consists of dislocation cells and cell-blocks (CBs) [20]. Within each CB, the deformation is accommodated by one set of slip systems that are different from those in neighboring CBs. The boundary between the neighboring CBs is called geometrically necessary boundary (GNB) since the boundary is necessary to accommodate the glide-induced lattice rotation in the adjoining CBs [20]. Figure 8 shows that GNBs may be parallel to the {111} slip planes or other crystallographic planes. The dislocation cell boundaries are called incidental boundaries since they are generated by the statistical mutual trappings of glide dislocations. Hansen and coworkers [16, 17, 19] observed two types of features on the GNBs in the rolling-induced microstructure. One is DDWs. The other is microbands (MBs) that are composed of small pancake-shaped cells (SPCs). The definition of MBs has been controversial and confusing since the nomenclature has been used to describe the long, thin plate-like regions observed in deformed copper [37–39]. The