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Severe plastic deformation (SPD) of titanium at near-ambient temperature
Acta Materialia 54 (2006) 3691–3700 www.actamat-journals.com
Severe plastic deformation (SPD) of titanium at near-ambient temperature M. Ravi Shankar a, Balkrishna C. Rao a, Seongeyl Lee a, Srinivasan Chandrasekar Alexander H. King b, W. Dale Compton a a
a,*
,
Center for Materials Processing and Tribology, School of Industrial Engineering, Purdue University, IE, GRIS, Purdue University 315 North Grant, West Lafayette, IN 47907-2023, USA b School of Materials Engineering, Purdue University, West Lafayette, IN 47907, USA Received 15 January 2006; received in revised form 29 March 2006; accepted 30 March 2006 Available online 14 June 2006
Abstract Severe plastic deformation of titanium at near-ambient temperature has been realized using large-strain machining. The material that received the highest levels of strain, the chip, is found to have a nanocrystalline structure with 100 nm sized grains. This microstructure is finer than that observed in titanium following equal channel angular pressing at elevated temperatures. The machined surface of the workpiece is found to have a graded microstructure composed of ultrafine grains at and near the surface and a preponderance of twins further into the subsurface. The parameters of the large-strain deformation field, namely strain rate and strain, are obtained using an adaptation of a particle image velocimetry technique. The microstructure characteristics are shown to be a consequence of the deformation field parameters, which activate both twinning and dislocation-mediated mechanisms of plastic flow. Ó 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; Ultrafine-grained microstructure; Titanium; Nanostructure
1. Introduction Severe plastic deformation (SPD) has emerged as an attractive approach for the creation of bulk metals and alloys with ultrafine-grained (UFG) microstructure and enhanced strength and hardness, due to its inherent simplicity [1–9]. In moderate-to high-strength materials of limited ductility, such as titanium and Al-6061T6, conventional SPD techniques such as equal channel angular pressing (ECAP) can only be performed at elevated temperatures where the flow characteristics of the material (e.g., increased ductility, reduction in yield strength) facilitate large-strain deformation. A common feature of conventional SPD techniques is the use of multiple passes of deformation (as in ECAP) or inhomogeneous deformation (as in high-pressure torsion, HPT) to impose the large strains needed for refine*
Corresponding author. Tel.: +1 76 54 94 36 23; fax: +1 76 54 94 54 48. E-mail address: [email protected] (S. Chandrasekar).
ment of the microstructure [10]. Recently, chip formation by machining has been shown to be an effective way to achieve SPD at near-ambient temperature, even for materials of high strength and limited ductility [8,9]. The large-strain deformation during chip formation under plane-strain conditions (Fig. 1) is somewhat akin to that in ECAP, although significantly larger strains can be imposed in a single pass of the cutting tool. The deformation strain in machining is controllable through an appropriate choice of the tool rake angle (a in Fig. 1) and can be varied easily in the range of 1–10 [8]. Thus chip formation by plane-strain machining can be viewed as a controlled large-strain deformation process; we call this large-strain machining. Interest in titanium and titanium alloys with UFG microstructures is motivated by their potential attractiveness for use as high-temperature structural materials. UFG titanium has been produced by deforming titanium to large values of strain using ECAP at temperatures of
1359-6454/$30.00 Ó 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.03.056
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shear in a narrow deformation zone, the ‘‘shear plane’’ in Fig. 1 [11]. Most of the grain refinement associated with the formation of the UFG chips has been attributed to the large shear strains imposed in this deformation zone [9]. The geometry of this deformation field and the associated shear strain (c) are determined by the shear plane angle (/) and the rake angle (a). These may be estimated using the equation [11] cos a c¼ ð1Þ sin / cosð/ aÞ where /, the orientation of the shear plane, is found from measurement of a0 and ac such that tan / ¼ Fig. 1. Schematic of large-strain machining.
400 °C [2–5]. In some of these studies, the microstructure refinement was found to be insufficient to produce the desired properties, and an additional deformation process, such as cold rolling, was used to further refine the microstructure [2,3,5]. Here, we demonstrate the successful adaptation of chip formation during large-strain machining, as an SPD process at near-ambient temperature for commercially pure titanium. The objectives are to elucidate the large-strain deformation behavior of titanium and to explore the properties of titanium with UFG microstructures. 1.1. Large-strain deformation in machining Fig. 1 shows a schematic of large-strain machining. Chip formation is often idealized as occurring by concentrated
a0 ac
1
cos a a0 ac
ð2Þ
sin a
While the deformation zone can be idealized for analytic simplicity as a shear plane, a zone of infinitesimal thickness, it is in reality a zone of finite dimensions. Deformation occurs typically over a fan-shaped region that extends ahead of the cutting tool and into the bulk material as illustrated in Fig. 1. Any volume of the bulk material that traverses this deformation zone and becomes part of the chip undergoes significant refinement of its microstructure [9,12]; the chip thus created by large-strain deformation is found to be nanostructured for a variety of materials [13]. 2. Experimental Commercially pure titanium was deformed to large strains at near-ambient temperature in a specially devised large-strain machining setup (Fig. 2). The titanium sample, in the form of a plate (bulk), had an initial grain size of
cutting tool
t
Chip Tool
Bulk
workpiece
t+∠ t Tool side view of the deformation field
Fig. 2. Schematic of the experimental setup used to acquire high-speed images of the deformation in machining. The images are captured with a highspeed CCD camera and analyzed using PIV to characterize the strain rates and strains associated with the deformation field.
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60 lm and a Vickers hardness of 144 kg/mm2 (300 g load). Characterization of the bulk titanium using X-ray diffraction confirmed the presence of only the hexagonal close-packed phase and an absence of any retained bodycentered cubic titanium. High-speed steel tools of rake angles +20° and 20° were then used to impose different levels of plastic strain in the chips. The typical width of the chips was 3 mm and the undeformed chip thickness (a0) value used in the experiment was 300 lm. A cutting speed of 10 mm/s, sufficiently low to ensure minimal temperature rise in the deformation zone, was employed while realizing the large-strain deformation. Fig. 2 shows a schematic of the experimental arrangement used to observe the deformation in large-strain machining. One side of the plate sample (workpiece) was constrained by a transparent glass block to ensure minimal side flow of the material during chip formation while enabling direct observation of material flow through the deformation zone. A sequence of images, in time steps (Dt) corresponding to the framing rate of a high-speed CCD camera (Kodak Motion Coder Analyzer Sr-Ultra), was collected to characterize the deformation field parameters, i.e., velocity, strain rate and strain. By following the motion of asperities in the material as recorded in the side view and applying an adaptation of a particle image velocimetry (PIV) technique [14], an image correlation method, the velocity field in the deformation zone was obtained [12]. The strain rate field was derived by differentiating the velocity field. Then, by integrating the strain rate field along particle trajectories through the deformation zone, the cumulative strain imposed in a volume element of the
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material as it traversed the deformation zone was evaluated. A particle trajectory could traverse the deformation zone and end up in the chip or in the surface/subsurface of the residual workpiece underneath the tool. Thus by appropriate selection of the particle trajectories, the strain at different locations in the deformation zone, chip and the subsurface of the workpiece could be determined. Details of this PIV technique for the assessment of the deformation field in machining can be found in Ref. [12]. Partially detached chip specimens similar to that shown in Fig. 3 were prepared by conventional metallography and observed under an optical microscope to study the microstructural features of the deformation. By superimposing the deformation field parameters onto the microstructure features, correlations could be established between the two. Hardness measurements were made using Vickers indentation at different locations in the chip, deformation zone and the bulk. Transmission electron microscopy (TEM) was used to study the internal microstructure of the chips created by the large-strain machining, allowing us to measure the grain sizes and observe the dislocation structures. For this purpose, electron-transparent chip specimens were prepared using a combination of mechanical polishing and electrolytic thinning. The chips were first polished down to a thickness of 100 lm on an abrasive polishing wheel. Disks of 3 mm in diameter were then punched out of the polished specimens. The disk specimens were made electron transparent by electrolytic jet-thinning (Struers Tenupol-5) using a solution of 60% methanol, 35% butanol and 5% perchloric acid at 30 °C and at 40 V. The electron
Fig. 3. Optical micrograph of metallographic features associated with the deformation and chip formation in large-strain machining of titanium with a +20° rake angle tool. Two distinct zones of deformation are highlighted. The fan-shaped deformation zone is characterized by a proliferation of twinning, elongated grain structures and retention of some of the metallographic features of the bulk. The zone of localized severe deformation may be identified with the onset of ‘‘destruction’’ of the metallographic features.
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transparent specimens were then studied using a JEOL 2000FX TEM instrument operating at 200 kV. 3. Results Figs. 3 and 4 (higher magnification image) show optical micrographs of a partially detached chip specimen machined from the titanium plate sample with a +20° rake angle tool. The equiaxed, coarse-grained microstructure of the bulk evolves, following deformation, into a metallographically ‘‘featureless’’ chip, as observed by optical microscopy. Such a transformation of the microstructure is characteristic of formation of UFG chips by large-strain deformation [9]. The evolution of the microstructure is somewhat gradual in a region ahead of the tool, labeled ‘‘fan-shaped deformation zone’’ in Fig. 3. This zone is characterized by the elongation of hitherto equiaxed grains in the bulk, due to shear deformation, but the sheared grains can still be individually resolved in this zone in Figs. 3 and 4. Subsequently, the microstructure evolves rather abruptly over a narrow region with the individual grains being rendered indistinguishable in the chip, a consequence of SPD occurring in this region. This narrow region is, hence, labeled as the zone of localized severe deformation (Fig. 3). It is expected that the largest increment of plastic strain is imposed in this zone with concomitant significant refinement of the microstructure, resulting in the apparently featureless chip. In Figs. 3 and 4, except for a few dispersed annealing twins, the undeformed bulk material is relatively free of twins. In the fan-shaped deformation zone, however, a greater density of twins can be seen (highlighted by the arrows in
Fig. 4). One of the grains circumscribed by a parallelogram in Fig. 4 contains a number of micro-twins that traverse the entire width of the grain. Fig. 4 highlights two specific grains circumscribed by parallelograms. Assuming that these grains were initially equiaxed in the bulk and could be circumscribed by rectangles (the parallelograms represent the now-sheared rectangles), the orientation of the parallelograms gives an indication of the direction and level of shear deformation in the material. The angle (h) through which the grains are sheared, h = 42°, corresponds to an effective strain value (e) given by tan h e ¼ pffiffiffi 0:52 3 Fig. 5 shows an optical micrograph of a cross-section through the machined surface and bulk subsurface left behind after chip formation, in the wake of the +20° rake angle tool. The zone closest to the machined surface is metallographically featureless, like the chip, typical of a material subjected to large values of plastic strain. Further into the subsurface, we note significant twinning. These metallographic characteristics are similar to those observed in the deformation zones ahead of the tool in Figs. 3 and 4. The machined subsurface can be visualized as the microstructure in the wake of chip formation that inherits its microstructural characteristics from these deformation zones. In Fig. 5, the zone closest to the machined surface is inherited from the metallographically featureless zone of localized severe deformation while the twinned subsurface is an offshoot of the fan-shaped deformation zone (also characterized by considerable twinning).
Fig. 4. Close-up of the deformation zone in Fig. 3. Parallelograms have been circumscribed around two of the sheared grains to characterize the shear strain. The arrows indicate significant twinning in the fan-shaped deformation zone. The solid white line demarcates the beginning of the zone of localized severe deformation and the destruction of metallographic features.
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Fig. 5. Optical micrograph of metallographic features on the machined surface and adjoining subsurface following material removal by chip formation with the +20° rake angle tool. The subsurface inherited from the fan-shaped deformation zone is characterized by a proliferation of twinning while that inherited from the zone of localized severe deformation is metallographically ‘‘featureless’’.
Fig. 6 shows the variation in the shear strain rate in the deformation zones, determined using PIV, for machining with the +20° rake angle tool. The shear strain rate increases gradually in the fan-shaped deformation zone.
The similarity between the morphology of the fan-shaped deformation zone determined using PIV in Fig. 6 and that observed metallographically in Fig. 3 may be noted. Transiting through the fan-shaped deformation zone, significant
Fig. 6. Shear strain rate field in the deformation zones obtained using PIV for the +20° rake angle tool. The fan-shaped deformation zone is identified with smaller values of the strain rate. The zone of localized severe deformation is identified with the higher strain rates. Units of the strain rates shown are per second.
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shearing occurs over a narrow zone just ahead of the tool. This is revealed by the larger shear strain rates prevailing in the zone of localized severe deformation in Fig. 6. This zone of localized severe deformation has earlier been associated with the obliteration of metallographic features in Figs. 3 and 4. Tracking the material from the undeformed bulk, through the deformation zones and into the chip or the machined surface, the total strain imposed in different regions of the material during chip formation can be evaluated using the strain rate field [12]. Three distinct particle trajectories, shown in Fig. 7, were used to evaluate the total effective strain. Fig. 8 gives the total effective strain at various points along these trajectories; the locations of these points are highlighted in Fig. 7. The effective strain value is seen to increase gradually between points A and B in the fan-shaped deformation zone. This is consistent with the gradual evolution of the microstructure in this zone as seen in Figs. 3 and 4. However, between points B and C 0 , the effective strain value increases steeply, corresponding to the zone of localized severe deformation. A good correspondence can be noted between the strain value calculated using PIV (Fig. 8) and that determined by observing the shearing of individual grains (e 0.52) using metallography for the locations of the sheared parallelograms in Fig. 4 that lie about midway between points B and C 0 in Fig. 7. It may be noted that paths 2 and 3 in Fig. 8 correspond to the plastic strains in the machined subsurface. Large values of strain are found closest to the machined surface (path-2) and the strain values decline rapidly further into the subsurface (path-3). Fig. 9 shows a TEM micrograph of a chip machined with the +20° rake angle tool. The effective strain imposed in this chip is about 1.4 as seen from Fig. 8. The microstructure
Path1
Path2
Path3
1.6 1.4
Total effective strain
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D' 1.2
C'
1 0.8
A
B
C
D
0.6 0.4 0.2 0
Distance along ABCD Fig. 8. Total effective strain accumulated along the various trajectories. The trajectories and points are identified in Fig. 7. The standard deviations estimated for the total effective strain are: Path 1, Point D 0 : 0.020; Path 2, Point D: 0.046; Path 3, Point D: 0.036.
consists of equiaxed, sub-100 nm grains interspersed with elongated, less developed subgrain structures. The Vickers hardness value of this chip was 230 kg/mm2, which is significantly greater than that of the bulk titanium (144 kg/ mm2) and consistent with its UFG microstructure. Fig. 10 shows a nanocrystalline microstructure resulting from deformation of the titanium to a shear strain (c) of 6 (effective strain 3.5) using a 20° rake angle tool; this level of deformation is much greater than that realized with the +20° rake angle tool. This shear strain value was calculated using the shear plane model (Eqs. (1) and (2)), since it was difficult to characterize the deformation field for this condition using PIV. The microstructure in Fig. 10 shows a greater level of refinement than that of Fig. 9. We find that
Fig. 7. Particle trajectories used to determine the total effective strain accumulated in different regions of titanium for the +20° rake angle tool. The total strain accumulated is obtained by integrating of the strain rate field along these particle trajectories. The intersection of the vertical white lines with the trajectories gives the location of the labeled points.
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Fig. 9. Bright-field TEM image of titanium chip created with an effective strain of 1.4 (+20° rake angle tool). The inset shows the corresponding SAD pattern. The arrow indicates a BCI pattern that is typically observed in titanium subjected to SPD at room temperature. The strain rate field and strains corresponding to this deformation can be seen in Figs. 6–8.
Fig. 10. Bright-field TEM image of titanium chip created with an effective strain of 3.5 (20° rake angle tool). The inset shows the corresponding SAD pattern.
most of the subgrain structures are well delineated in Fig. 10 corresponding to a material that has undergone considerable microstructure refinement, due to the larger values of shear strain. Comparing the selected area diffraction (SAD) patterns in Figs. 9 and 10, it is also apparent that the chip in Fig. 10 consists of grain boundaries of more significant misorientation; this is consistent again with the lar-
ger levels of imposed deformation. While the SAD patterns in themselves are not sufficient to indicate the misorientation values, they can still be used to judge qualitatively that the misorientations corresponding to the ring-like pattern in Fig. 10 are larger than those that produce the smeared single-crystal pattern in Fig. 9. The Vickers hardness of this nanocrystalline chip was 247 kg/mm2.
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4. Discussion A study has been made of the deformation of titanium in large-strain machining. The experimental observations have demonstrated that SPD of titanium at near-ambient temperature can be realized using this technique. The chips created by SPD have been shown to be composed of UFG microstructures, with those produced at the highest levels of strain having 100 nm sized grains. A graded microstructure has been noted at and near the surface of the residual workpiece. These observations can be rationalized based on consideration of the deformation field parameters and the deformation temperature. 4.1. Deformation field Two distinct deformation zones, each characterized by different levels of strain and metallographic features, have been found to constitute the deformation field in large-strain machining. In a material like titanium, these characteristics lead to some interesting microstructural consequences. At low strains, twinning can contribute significantly to accommodation of the deformation [6,15,16]. The microstructure of titanium subjected to a single pass of ECAP typically consists of a number of deformation twins [6]. In large-strain machining, the analog of the deformation occurring in a single pass of ECAP is that prevailing in the fan-shaped deformation zone; this zone is characterized by a monotonic and a gradual increase in strain to relatively low values (e 0.52). As seen in Figs. 6 and 8, when the material enters the fan-shaped deformation zone, it initially encounters relatively small values of strain rate resulting in small accumulated plastic strains. Such small strains can be accommodated through twinning [15]. Upon undergoing further deformation in the fanshaped deformation zone, the microstructure consists of elongated grain structures and twins, the twins being those formed during the initial stages of the deformation (Figs. 3 and 4). Adjoining the fan-shaped zone along its upper boundary in Fig. 3 is a zone of localized severe deformation wherein there is a complete obliteration of the metallographic features. Based on observations of sheared grains in Fig. 4 near the upper boundary of the fan-shaped deformation zone, we estimated the effective strain in this region as 0.52. This value agrees well with that estimated using PIV at this location. A number of twins can also be seen within the deforming grains in this region. Twinning can accommodate only a portion of the plastic strains in the fan-shaped deformation zone. Even in this zone, a significant portion of the strains are expected to be accommodated by dislocation slip resulting in some refinement of the microstructure since the strains reach values as high as 0.5 close to the upper boundary of the fanshaped deformation zone. We were not able to resolve any twins in the TEM micrographs of the chips in Figs. 9 and 10. This is possibly a result of the large dislocation density making observation of the micro-twins difficult, rather
than evidence of absence of twins in the chips. However, it has also been suggested that in a material with a refined microstructure, even in one with limited slip systems such as titanium at near-ambient temperature, compatible deformation can occur through coordinated slip in the fine grains without any contribution from twinning [7]. It has been shown in large-strain machining that the microstructure can undergo significant refinement even before the material enters the zone of localized severe deformation [9]. Thus, deformation in the zone of localized severe deformation may simply be accommodated by collective, compatible deformation in the already refined microstructure. The coexistence of different modes of deformation is also reflected in the microstructure of the machined surface left behind in the wake of the tool following chip formation. As seen in Fig. 5, the zone nearest to the machined surface is largely featureless while the material further into the subsurface indicates considerable twinning. These observations can be rationalized using the strains calculated along two particle trajectories in the machined subsurface in Fig. 7. Fig. 8 (path-2) demonstrates the presence of a zone of material subjected to large strains close to the machined surface, thus confirming the genesis of the metallographically featureless surface observed in Fig. 5. The machined surface here is likely nanostructured, mirroring that of the chip. Further into the subsurface, along path-3 in Figs. 7 and 8, smaller values of strain prevail and in accord with the argument presented for strain accommodation in the fan-shaped deformation zone, we expect and, indeed, see significant twinning. We note that this graded microstructure is quite similar to that observed in titanium that has been subjected to surface mechanical attrition treatment (SMAT) [17]. Thus, large-strain machining not only offers a route for the manufacture of nanostructured titanium but also the creation of a surface with a graded microstructure – a microstructure characterized by 100 nm grains at and near the machined surface and a preponderance of twins further into the subsurface. Significant twinning in the early stages of deformation may also influence the micromechanics of deformation at a later stage such as in the zone of localized severe deformation. A large density of twins can hinder plastic flow by acting as barriers to slip and enhance dislocation pinning [18]. This effect can be similar to that of second-phase particles and aid in the refinement of microstructure during SPD [19]. 4.2. Effect of deformation temperature Deformation temperature can play a significant role in the formation of UFG microstructures by affecting the retention of dislocations and evolution of dislocation cells and subgrain structures. Usually, the lower the deformation temperature, the more refined is the microstructure, due to greater retention of the dislocations. In comparison, at higher deformation temperatures, thermally activated annihilation of dislocations leads to a coarser
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microstructure. Consequently, large-strain machining of titanium at near-ambient temperature results in a finer microstructure (Figs. 9 and 10) compared to ECAP at elevated temperatures [2–5]. These observations in titanium parallel those in Al-6061 wherein the microstructure of chips created by large-strain machining at near-ambient temperature was finer than that resulting from ECAP at elevated temperatures [8]. In a material like titanium, the deformation temperature can determine the characteristics of the chip material not only through modification of the rate of annihilation of the dislocations, but also by modifying the micromechanics of flow by altering the slip systems that are operative during deformation. It has been suggested that the preponderance of slip on the f0 0 0 1gh1 1 2 0i system can modify the characteristics of the microstructure resulting from ECAP of titanium at temperatures of 400 °C [4]. This slip system is usually associated with prolific cross-slip with activation energy of 2.5 eV [20]. SPD by ECAP at these higher temperatures, therefore, may be associated with inefficient pinning and storage of dislocations and, consequently, limited microstructure refinement. During SPD at temperatures of 400 °C, the recovery mechanism aided by cross-slip can also lead to clearly defined grain boundaries, though inevitably accompanied by a reduction in the dislocation density [4]. In contrast for titanium deformed to large strains at near-ambient temperature, such as by large-strain machining, dislocation storage is likely to be more efficient. However, the resulting subgrain boundaries are likely to be non-equilibrium structures similar to those observed in titanium subjected to ECAP at elevated temperature followed by cold rolling [5]. Fig. 9 also displays the characteristic banded contrast image (BCI) pattern that has been attributed to closely spaced parallel dislocations typically encountered in titanium deformed at low temperatures [5]. The production of chips from titanium and its alloys by large-strain machining offers opportunities to create structural materials with mechanical properties that are influenced by the UFG microstructure of the chips. Large-strain machining can be directly adapted to make nanostructured titanium-based materials (chips) in the form of foils, plates, wires and filaments. Alternatively, comminution (e.g., attrition milling, cryo-milling) of the nanostructured chips is a possible route to large-scale production of nanostructured particles, which can be consolidated and densified into bulk monolithic materials using appropriate consolidation processes. An important consideration, as in consolidating any nanostructured material, is to limit exposure to high temperatures in order to suppress grain growth. While none of these processes have been fully developed for titanium-based materials, preliminary results suggest that the opportunity exists to create advanced structural materials with new and interesting combinations of properties. Finally, the existence of a graded microstructure through the residual workpiece surface suggests possibilities for the engineering of surfaces with interesting tribological and fatigue properties.
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5. Conclusions SPD of titanium at near-ambient temperature has been demonstrated using large-strain machining. This process is shown to create a microstructure in the chips that is significantly finer than that observed in titanium produced by SPD via ECAP at elevated temperatures, thus indicating a strong effect of deformation temperature on the formation of the UFG material. The residual workpiece is characterized by a graded microstructure composed of ultrafine grains at and near the surface and a preponderance of twins further into the subsurface. These microstructural characteristics are shown to be a consequence of the large strain deformation field wherein both twinning and dislocation-mediated plasticity are active. The large-strain deformation field is composed of two distinct zones: a fan-shaped deformation zone that extends ahead of the cutting tool and is characterized by relatively small values of strain rate and accumulated plastic strain, and a zone of localized severe deformation with large strain rates and strains. A preponderance of twinning and elongated grain structures are seen in the fan-shaped deformation zone. Significant microstructure refinement occurs in the localized zone of severe deformation typified by loss of metallographic features. Concomitantly, slip is expected to be the dominant deformation mode in this zone. Likewise, closest to the machined surface, large values of strain and a metallographically featureless microstructure are observed. This microstructure is likely to mirror that of the nanostructured chip. Further into the subsurface, the values of strain are smaller and significant twinning is observed. Acknowledgements We thank the US Department of Energy (Grant 4000031768 via UT-Batelle), Ford Motor Company, Oak Ridge National Laboratory (ORNL), the NSF (Grant DMI 0500216) and the State of Indiana 21st Century Research and Technology Fund for support of this work. Additional thanks are also due to Drs. Ray Johnson (ORNL) and Andrew Sherman (Ford) for their encouragement of the studies. References [1] Ferrasse S, Segal VM, Hartwig KT, Goforth RE. J Mater Res 1997;12:1253–61. [2] Stolyarov VV, Zhu YT, Lowe TC, Valiev RZ. J Nanosci Nanotechnol 2001;1:237–42. [3] Stolyarov VV, Zhu YT, Alexandarov IV, Lowe TC, Valiev RZ. Mater Sci Eng A 2003;343:43–50. [4] Stolyarov VV, Zhu YT, Alexandarov IV, Lowe TC, Valiev RZ. Mater Sci Eng A 2001;299:59–67. [5] Zhu YT, Huang JY, Gubicza J, Ungar T, Wang YM, Ma E, et al. J Mater Res 2003;18:1908–17. [6] Kim I, Kim J, Shin DH, Liao XZ, Zhu YT. Scripta Mater 2003;48: 813–7.
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[7] Shin DH, Kim I, Kim J, Kim YS, Semiatin SL. Acta Mater 2003;51: 983–96. [8] Shankar MR, Chandrasekar S, King AH, Compton WD. Acta Mater 2005;53:4781–93. [9] Shankar MR, Chandrasekar S, Compton WD, King AH. Mater Sci Eng A 2005;410–411:364–8. [10] Furukawa M, Horita Z, Nemoto M, Langdon TG. Mater Sci Eng A 2002;324:82–9. [11] Shaw MC. Metal cutting principles. Oxford series on advanced manufacturing. Oxford: Clarendon Press; 1984. [12] Lee S, Hwang J, Shankar MR, Chandrasekar S, Compton WD. Metall Mater Trans A 2006;37(5):1633–43.
[13] Swaminathan S, Shankar MR, Lee S, Hwang J, King AH, Kezar RF, et al. Mater Sci Eng A 2005;410–411:358–63. [14] Adrian RJ. Laser Anemom Fluid Mech 1988;3:115–29. [15] Minonishi Y, Morozumi S, Yoshinaga H. Scripta Metall 1985;19: 1241–5. [16] Serra A, Pond RC, Bacon DJ. Acta Metall Mater 1991;39:1469–80. [17] Zhu KY, Vassel A, Brisset F, Lu K, Lu J. Acta Mater 2004;52: 4101–10. [18] Serra A, Bacon DJ, Pond RC. Metall Mater Trans A 2002;33: 809–12. [19] Apps PJ, Bowen JR, Prangnell PB. Acta Mater 2003;51:2811–22. [20] Akhtar A. Metall Trans A 1975;6:1105–13.
Severe plastic deformation (SPD) of titanium at near-ambient temperature M. Ravi Shankar a, Balkrishna C. Rao a, Seongeyl Lee a, Srinivasan Chandrasekar Alexander H. King b, W. Dale Compton a a
a,*
,
Center for Materials Processing and Tribology, School of Industrial Engineering, Purdue University, IE, GRIS, Purdue University 315 North Grant, West Lafayette, IN 47907-2023, USA b School of Materials Engineering, Purdue University, West Lafayette, IN 47907, USA Received 15 January 2006; received in revised form 29 March 2006; accepted 30 March 2006 Available online 14 June 2006
Abstract Severe plastic deformation of titanium at near-ambient temperature has been realized using large-strain machining. The material that received the highest levels of strain, the chip, is found to have a nanocrystalline structure with 100 nm sized grains. This microstructure is finer than that observed in titanium following equal channel angular pressing at elevated temperatures. The machined surface of the workpiece is found to have a graded microstructure composed of ultrafine grains at and near the surface and a preponderance of twins further into the subsurface. The parameters of the large-strain deformation field, namely strain rate and strain, are obtained using an adaptation of a particle image velocimetry technique. The microstructure characteristics are shown to be a consequence of the deformation field parameters, which activate both twinning and dislocation-mediated mechanisms of plastic flow. Ó 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; Ultrafine-grained microstructure; Titanium; Nanostructure
1. Introduction Severe plastic deformation (SPD) has emerged as an attractive approach for the creation of bulk metals and alloys with ultrafine-grained (UFG) microstructure and enhanced strength and hardness, due to its inherent simplicity [1–9]. In moderate-to high-strength materials of limited ductility, such as titanium and Al-6061T6, conventional SPD techniques such as equal channel angular pressing (ECAP) can only be performed at elevated temperatures where the flow characteristics of the material (e.g., increased ductility, reduction in yield strength) facilitate large-strain deformation. A common feature of conventional SPD techniques is the use of multiple passes of deformation (as in ECAP) or inhomogeneous deformation (as in high-pressure torsion, HPT) to impose the large strains needed for refine*
Corresponding author. Tel.: +1 76 54 94 36 23; fax: +1 76 54 94 54 48. E-mail address: [email protected] (S. Chandrasekar).
ment of the microstructure [10]. Recently, chip formation by machining has been shown to be an effective way to achieve SPD at near-ambient temperature, even for materials of high strength and limited ductility [8,9]. The large-strain deformation during chip formation under plane-strain conditions (Fig. 1) is somewhat akin to that in ECAP, although significantly larger strains can be imposed in a single pass of the cutting tool. The deformation strain in machining is controllable through an appropriate choice of the tool rake angle (a in Fig. 1) and can be varied easily in the range of 1–10 [8]. Thus chip formation by plane-strain machining can be viewed as a controlled large-strain deformation process; we call this large-strain machining. Interest in titanium and titanium alloys with UFG microstructures is motivated by their potential attractiveness for use as high-temperature structural materials. UFG titanium has been produced by deforming titanium to large values of strain using ECAP at temperatures of
1359-6454/$30.00 Ó 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.03.056
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shear in a narrow deformation zone, the ‘‘shear plane’’ in Fig. 1 [11]. Most of the grain refinement associated with the formation of the UFG chips has been attributed to the large shear strains imposed in this deformation zone [9]. The geometry of this deformation field and the associated shear strain (c) are determined by the shear plane angle (/) and the rake angle (a). These may be estimated using the equation [11] cos a c¼ ð1Þ sin / cosð/ aÞ where /, the orientation of the shear plane, is found from measurement of a0 and ac such that tan / ¼ Fig. 1. Schematic of large-strain machining.
400 °C [2–5]. In some of these studies, the microstructure refinement was found to be insufficient to produce the desired properties, and an additional deformation process, such as cold rolling, was used to further refine the microstructure [2,3,5]. Here, we demonstrate the successful adaptation of chip formation during large-strain machining, as an SPD process at near-ambient temperature for commercially pure titanium. The objectives are to elucidate the large-strain deformation behavior of titanium and to explore the properties of titanium with UFG microstructures. 1.1. Large-strain deformation in machining Fig. 1 shows a schematic of large-strain machining. Chip formation is often idealized as occurring by concentrated
a0 ac
1
cos a a0 ac
ð2Þ
sin a
While the deformation zone can be idealized for analytic simplicity as a shear plane, a zone of infinitesimal thickness, it is in reality a zone of finite dimensions. Deformation occurs typically over a fan-shaped region that extends ahead of the cutting tool and into the bulk material as illustrated in Fig. 1. Any volume of the bulk material that traverses this deformation zone and becomes part of the chip undergoes significant refinement of its microstructure [9,12]; the chip thus created by large-strain deformation is found to be nanostructured for a variety of materials [13]. 2. Experimental Commercially pure titanium was deformed to large strains at near-ambient temperature in a specially devised large-strain machining setup (Fig. 2). The titanium sample, in the form of a plate (bulk), had an initial grain size of
cutting tool
t
Chip Tool
Bulk
workpiece
t+∠ t Tool side view of the deformation field
Fig. 2. Schematic of the experimental setup used to acquire high-speed images of the deformation in machining. The images are captured with a highspeed CCD camera and analyzed using PIV to characterize the strain rates and strains associated with the deformation field.
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60 lm and a Vickers hardness of 144 kg/mm2 (300 g load). Characterization of the bulk titanium using X-ray diffraction confirmed the presence of only the hexagonal close-packed phase and an absence of any retained bodycentered cubic titanium. High-speed steel tools of rake angles +20° and 20° were then used to impose different levels of plastic strain in the chips. The typical width of the chips was 3 mm and the undeformed chip thickness (a0) value used in the experiment was 300 lm. A cutting speed of 10 mm/s, sufficiently low to ensure minimal temperature rise in the deformation zone, was employed while realizing the large-strain deformation. Fig. 2 shows a schematic of the experimental arrangement used to observe the deformation in large-strain machining. One side of the plate sample (workpiece) was constrained by a transparent glass block to ensure minimal side flow of the material during chip formation while enabling direct observation of material flow through the deformation zone. A sequence of images, in time steps (Dt) corresponding to the framing rate of a high-speed CCD camera (Kodak Motion Coder Analyzer Sr-Ultra), was collected to characterize the deformation field parameters, i.e., velocity, strain rate and strain. By following the motion of asperities in the material as recorded in the side view and applying an adaptation of a particle image velocimetry (PIV) technique [14], an image correlation method, the velocity field in the deformation zone was obtained [12]. The strain rate field was derived by differentiating the velocity field. Then, by integrating the strain rate field along particle trajectories through the deformation zone, the cumulative strain imposed in a volume element of the
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material as it traversed the deformation zone was evaluated. A particle trajectory could traverse the deformation zone and end up in the chip or in the surface/subsurface of the residual workpiece underneath the tool. Thus by appropriate selection of the particle trajectories, the strain at different locations in the deformation zone, chip and the subsurface of the workpiece could be determined. Details of this PIV technique for the assessment of the deformation field in machining can be found in Ref. [12]. Partially detached chip specimens similar to that shown in Fig. 3 were prepared by conventional metallography and observed under an optical microscope to study the microstructural features of the deformation. By superimposing the deformation field parameters onto the microstructure features, correlations could be established between the two. Hardness measurements were made using Vickers indentation at different locations in the chip, deformation zone and the bulk. Transmission electron microscopy (TEM) was used to study the internal microstructure of the chips created by the large-strain machining, allowing us to measure the grain sizes and observe the dislocation structures. For this purpose, electron-transparent chip specimens were prepared using a combination of mechanical polishing and electrolytic thinning. The chips were first polished down to a thickness of 100 lm on an abrasive polishing wheel. Disks of 3 mm in diameter were then punched out of the polished specimens. The disk specimens were made electron transparent by electrolytic jet-thinning (Struers Tenupol-5) using a solution of 60% methanol, 35% butanol and 5% perchloric acid at 30 °C and at 40 V. The electron
Fig. 3. Optical micrograph of metallographic features associated with the deformation and chip formation in large-strain machining of titanium with a +20° rake angle tool. Two distinct zones of deformation are highlighted. The fan-shaped deformation zone is characterized by a proliferation of twinning, elongated grain structures and retention of some of the metallographic features of the bulk. The zone of localized severe deformation may be identified with the onset of ‘‘destruction’’ of the metallographic features.
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transparent specimens were then studied using a JEOL 2000FX TEM instrument operating at 200 kV. 3. Results Figs. 3 and 4 (higher magnification image) show optical micrographs of a partially detached chip specimen machined from the titanium plate sample with a +20° rake angle tool. The equiaxed, coarse-grained microstructure of the bulk evolves, following deformation, into a metallographically ‘‘featureless’’ chip, as observed by optical microscopy. Such a transformation of the microstructure is characteristic of formation of UFG chips by large-strain deformation [9]. The evolution of the microstructure is somewhat gradual in a region ahead of the tool, labeled ‘‘fan-shaped deformation zone’’ in Fig. 3. This zone is characterized by the elongation of hitherto equiaxed grains in the bulk, due to shear deformation, but the sheared grains can still be individually resolved in this zone in Figs. 3 and 4. Subsequently, the microstructure evolves rather abruptly over a narrow region with the individual grains being rendered indistinguishable in the chip, a consequence of SPD occurring in this region. This narrow region is, hence, labeled as the zone of localized severe deformation (Fig. 3). It is expected that the largest increment of plastic strain is imposed in this zone with concomitant significant refinement of the microstructure, resulting in the apparently featureless chip. In Figs. 3 and 4, except for a few dispersed annealing twins, the undeformed bulk material is relatively free of twins. In the fan-shaped deformation zone, however, a greater density of twins can be seen (highlighted by the arrows in
Fig. 4). One of the grains circumscribed by a parallelogram in Fig. 4 contains a number of micro-twins that traverse the entire width of the grain. Fig. 4 highlights two specific grains circumscribed by parallelograms. Assuming that these grains were initially equiaxed in the bulk and could be circumscribed by rectangles (the parallelograms represent the now-sheared rectangles), the orientation of the parallelograms gives an indication of the direction and level of shear deformation in the material. The angle (h) through which the grains are sheared, h = 42°, corresponds to an effective strain value (e) given by tan h e ¼ pffiffiffi 0:52 3 Fig. 5 shows an optical micrograph of a cross-section through the machined surface and bulk subsurface left behind after chip formation, in the wake of the +20° rake angle tool. The zone closest to the machined surface is metallographically featureless, like the chip, typical of a material subjected to large values of plastic strain. Further into the subsurface, we note significant twinning. These metallographic characteristics are similar to those observed in the deformation zones ahead of the tool in Figs. 3 and 4. The machined subsurface can be visualized as the microstructure in the wake of chip formation that inherits its microstructural characteristics from these deformation zones. In Fig. 5, the zone closest to the machined surface is inherited from the metallographically featureless zone of localized severe deformation while the twinned subsurface is an offshoot of the fan-shaped deformation zone (also characterized by considerable twinning).
Fig. 4. Close-up of the deformation zone in Fig. 3. Parallelograms have been circumscribed around two of the sheared grains to characterize the shear strain. The arrows indicate significant twinning in the fan-shaped deformation zone. The solid white line demarcates the beginning of the zone of localized severe deformation and the destruction of metallographic features.
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Fig. 5. Optical micrograph of metallographic features on the machined surface and adjoining subsurface following material removal by chip formation with the +20° rake angle tool. The subsurface inherited from the fan-shaped deformation zone is characterized by a proliferation of twinning while that inherited from the zone of localized severe deformation is metallographically ‘‘featureless’’.
Fig. 6 shows the variation in the shear strain rate in the deformation zones, determined using PIV, for machining with the +20° rake angle tool. The shear strain rate increases gradually in the fan-shaped deformation zone.
The similarity between the morphology of the fan-shaped deformation zone determined using PIV in Fig. 6 and that observed metallographically in Fig. 3 may be noted. Transiting through the fan-shaped deformation zone, significant
Fig. 6. Shear strain rate field in the deformation zones obtained using PIV for the +20° rake angle tool. The fan-shaped deformation zone is identified with smaller values of the strain rate. The zone of localized severe deformation is identified with the higher strain rates. Units of the strain rates shown are per second.
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shearing occurs over a narrow zone just ahead of the tool. This is revealed by the larger shear strain rates prevailing in the zone of localized severe deformation in Fig. 6. This zone of localized severe deformation has earlier been associated with the obliteration of metallographic features in Figs. 3 and 4. Tracking the material from the undeformed bulk, through the deformation zones and into the chip or the machined surface, the total strain imposed in different regions of the material during chip formation can be evaluated using the strain rate field [12]. Three distinct particle trajectories, shown in Fig. 7, were used to evaluate the total effective strain. Fig. 8 gives the total effective strain at various points along these trajectories; the locations of these points are highlighted in Fig. 7. The effective strain value is seen to increase gradually between points A and B in the fan-shaped deformation zone. This is consistent with the gradual evolution of the microstructure in this zone as seen in Figs. 3 and 4. However, between points B and C 0 , the effective strain value increases steeply, corresponding to the zone of localized severe deformation. A good correspondence can be noted between the strain value calculated using PIV (Fig. 8) and that determined by observing the shearing of individual grains (e 0.52) using metallography for the locations of the sheared parallelograms in Fig. 4 that lie about midway between points B and C 0 in Fig. 7. It may be noted that paths 2 and 3 in Fig. 8 correspond to the plastic strains in the machined subsurface. Large values of strain are found closest to the machined surface (path-2) and the strain values decline rapidly further into the subsurface (path-3). Fig. 9 shows a TEM micrograph of a chip machined with the +20° rake angle tool. The effective strain imposed in this chip is about 1.4 as seen from Fig. 8. The microstructure
Path1
Path2
Path3
1.6 1.4
Total effective strain
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D' 1.2
C'
1 0.8
A
B
C
D
0.6 0.4 0.2 0
Distance along ABCD Fig. 8. Total effective strain accumulated along the various trajectories. The trajectories and points are identified in Fig. 7. The standard deviations estimated for the total effective strain are: Path 1, Point D 0 : 0.020; Path 2, Point D: 0.046; Path 3, Point D: 0.036.
consists of equiaxed, sub-100 nm grains interspersed with elongated, less developed subgrain structures. The Vickers hardness value of this chip was 230 kg/mm2, which is significantly greater than that of the bulk titanium (144 kg/ mm2) and consistent with its UFG microstructure. Fig. 10 shows a nanocrystalline microstructure resulting from deformation of the titanium to a shear strain (c) of 6 (effective strain 3.5) using a 20° rake angle tool; this level of deformation is much greater than that realized with the +20° rake angle tool. This shear strain value was calculated using the shear plane model (Eqs. (1) and (2)), since it was difficult to characterize the deformation field for this condition using PIV. The microstructure in Fig. 10 shows a greater level of refinement than that of Fig. 9. We find that
Fig. 7. Particle trajectories used to determine the total effective strain accumulated in different regions of titanium for the +20° rake angle tool. The total strain accumulated is obtained by integrating of the strain rate field along these particle trajectories. The intersection of the vertical white lines with the trajectories gives the location of the labeled points.
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Fig. 9. Bright-field TEM image of titanium chip created with an effective strain of 1.4 (+20° rake angle tool). The inset shows the corresponding SAD pattern. The arrow indicates a BCI pattern that is typically observed in titanium subjected to SPD at room temperature. The strain rate field and strains corresponding to this deformation can be seen in Figs. 6–8.
Fig. 10. Bright-field TEM image of titanium chip created with an effective strain of 3.5 (20° rake angle tool). The inset shows the corresponding SAD pattern.
most of the subgrain structures are well delineated in Fig. 10 corresponding to a material that has undergone considerable microstructure refinement, due to the larger values of shear strain. Comparing the selected area diffraction (SAD) patterns in Figs. 9 and 10, it is also apparent that the chip in Fig. 10 consists of grain boundaries of more significant misorientation; this is consistent again with the lar-
ger levels of imposed deformation. While the SAD patterns in themselves are not sufficient to indicate the misorientation values, they can still be used to judge qualitatively that the misorientations corresponding to the ring-like pattern in Fig. 10 are larger than those that produce the smeared single-crystal pattern in Fig. 9. The Vickers hardness of this nanocrystalline chip was 247 kg/mm2.
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4. Discussion A study has been made of the deformation of titanium in large-strain machining. The experimental observations have demonstrated that SPD of titanium at near-ambient temperature can be realized using this technique. The chips created by SPD have been shown to be composed of UFG microstructures, with those produced at the highest levels of strain having 100 nm sized grains. A graded microstructure has been noted at and near the surface of the residual workpiece. These observations can be rationalized based on consideration of the deformation field parameters and the deformation temperature. 4.1. Deformation field Two distinct deformation zones, each characterized by different levels of strain and metallographic features, have been found to constitute the deformation field in large-strain machining. In a material like titanium, these characteristics lead to some interesting microstructural consequences. At low strains, twinning can contribute significantly to accommodation of the deformation [6,15,16]. The microstructure of titanium subjected to a single pass of ECAP typically consists of a number of deformation twins [6]. In large-strain machining, the analog of the deformation occurring in a single pass of ECAP is that prevailing in the fan-shaped deformation zone; this zone is characterized by a monotonic and a gradual increase in strain to relatively low values (e 0.52). As seen in Figs. 6 and 8, when the material enters the fan-shaped deformation zone, it initially encounters relatively small values of strain rate resulting in small accumulated plastic strains. Such small strains can be accommodated through twinning [15]. Upon undergoing further deformation in the fanshaped deformation zone, the microstructure consists of elongated grain structures and twins, the twins being those formed during the initial stages of the deformation (Figs. 3 and 4). Adjoining the fan-shaped zone along its upper boundary in Fig. 3 is a zone of localized severe deformation wherein there is a complete obliteration of the metallographic features. Based on observations of sheared grains in Fig. 4 near the upper boundary of the fan-shaped deformation zone, we estimated the effective strain in this region as 0.52. This value agrees well with that estimated using PIV at this location. A number of twins can also be seen within the deforming grains in this region. Twinning can accommodate only a portion of the plastic strains in the fan-shaped deformation zone. Even in this zone, a significant portion of the strains are expected to be accommodated by dislocation slip resulting in some refinement of the microstructure since the strains reach values as high as 0.5 close to the upper boundary of the fanshaped deformation zone. We were not able to resolve any twins in the TEM micrographs of the chips in Figs. 9 and 10. This is possibly a result of the large dislocation density making observation of the micro-twins difficult, rather
than evidence of absence of twins in the chips. However, it has also been suggested that in a material with a refined microstructure, even in one with limited slip systems such as titanium at near-ambient temperature, compatible deformation can occur through coordinated slip in the fine grains without any contribution from twinning [7]. It has been shown in large-strain machining that the microstructure can undergo significant refinement even before the material enters the zone of localized severe deformation [9]. Thus, deformation in the zone of localized severe deformation may simply be accommodated by collective, compatible deformation in the already refined microstructure. The coexistence of different modes of deformation is also reflected in the microstructure of the machined surface left behind in the wake of the tool following chip formation. As seen in Fig. 5, the zone nearest to the machined surface is largely featureless while the material further into the subsurface indicates considerable twinning. These observations can be rationalized using the strains calculated along two particle trajectories in the machined subsurface in Fig. 7. Fig. 8 (path-2) demonstrates the presence of a zone of material subjected to large strains close to the machined surface, thus confirming the genesis of the metallographically featureless surface observed in Fig. 5. The machined surface here is likely nanostructured, mirroring that of the chip. Further into the subsurface, along path-3 in Figs. 7 and 8, smaller values of strain prevail and in accord with the argument presented for strain accommodation in the fan-shaped deformation zone, we expect and, indeed, see significant twinning. We note that this graded microstructure is quite similar to that observed in titanium that has been subjected to surface mechanical attrition treatment (SMAT) [17]. Thus, large-strain machining not only offers a route for the manufacture of nanostructured titanium but also the creation of a surface with a graded microstructure – a microstructure characterized by 100 nm grains at and near the machined surface and a preponderance of twins further into the subsurface. Significant twinning in the early stages of deformation may also influence the micromechanics of deformation at a later stage such as in the zone of localized severe deformation. A large density of twins can hinder plastic flow by acting as barriers to slip and enhance dislocation pinning [18]. This effect can be similar to that of second-phase particles and aid in the refinement of microstructure during SPD [19]. 4.2. Effect of deformation temperature Deformation temperature can play a significant role in the formation of UFG microstructures by affecting the retention of dislocations and evolution of dislocation cells and subgrain structures. Usually, the lower the deformation temperature, the more refined is the microstructure, due to greater retention of the dislocations. In comparison, at higher deformation temperatures, thermally activated annihilation of dislocations leads to a coarser
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microstructure. Consequently, large-strain machining of titanium at near-ambient temperature results in a finer microstructure (Figs. 9 and 10) compared to ECAP at elevated temperatures [2–5]. These observations in titanium parallel those in Al-6061 wherein the microstructure of chips created by large-strain machining at near-ambient temperature was finer than that resulting from ECAP at elevated temperatures [8]. In a material like titanium, the deformation temperature can determine the characteristics of the chip material not only through modification of the rate of annihilation of the dislocations, but also by modifying the micromechanics of flow by altering the slip systems that are operative during deformation. It has been suggested that the preponderance of slip on the f0 0 0 1gh1 1 2 0i system can modify the characteristics of the microstructure resulting from ECAP of titanium at temperatures of 400 °C [4]. This slip system is usually associated with prolific cross-slip with activation energy of 2.5 eV [20]. SPD by ECAP at these higher temperatures, therefore, may be associated with inefficient pinning and storage of dislocations and, consequently, limited microstructure refinement. During SPD at temperatures of 400 °C, the recovery mechanism aided by cross-slip can also lead to clearly defined grain boundaries, though inevitably accompanied by a reduction in the dislocation density [4]. In contrast for titanium deformed to large strains at near-ambient temperature, such as by large-strain machining, dislocation storage is likely to be more efficient. However, the resulting subgrain boundaries are likely to be non-equilibrium structures similar to those observed in titanium subjected to ECAP at elevated temperature followed by cold rolling [5]. Fig. 9 also displays the characteristic banded contrast image (BCI) pattern that has been attributed to closely spaced parallel dislocations typically encountered in titanium deformed at low temperatures [5]. The production of chips from titanium and its alloys by large-strain machining offers opportunities to create structural materials with mechanical properties that are influenced by the UFG microstructure of the chips. Large-strain machining can be directly adapted to make nanostructured titanium-based materials (chips) in the form of foils, plates, wires and filaments. Alternatively, comminution (e.g., attrition milling, cryo-milling) of the nanostructured chips is a possible route to large-scale production of nanostructured particles, which can be consolidated and densified into bulk monolithic materials using appropriate consolidation processes. An important consideration, as in consolidating any nanostructured material, is to limit exposure to high temperatures in order to suppress grain growth. While none of these processes have been fully developed for titanium-based materials, preliminary results suggest that the opportunity exists to create advanced structural materials with new and interesting combinations of properties. Finally, the existence of a graded microstructure through the residual workpiece surface suggests possibilities for the engineering of surfaces with interesting tribological and fatigue properties.
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5. Conclusions SPD of titanium at near-ambient temperature has been demonstrated using large-strain machining. This process is shown to create a microstructure in the chips that is significantly finer than that observed in titanium produced by SPD via ECAP at elevated temperatures, thus indicating a strong effect of deformation temperature on the formation of the UFG material. The residual workpiece is characterized by a graded microstructure composed of ultrafine grains at and near the surface and a preponderance of twins further into the subsurface. These microstructural characteristics are shown to be a consequence of the large strain deformation field wherein both twinning and dislocation-mediated plasticity are active. The large-strain deformation field is composed of two distinct zones: a fan-shaped deformation zone that extends ahead of the cutting tool and is characterized by relatively small values of strain rate and accumulated plastic strain, and a zone of localized severe deformation with large strain rates and strains. A preponderance of twinning and elongated grain structures are seen in the fan-shaped deformation zone. Significant microstructure refinement occurs in the localized zone of severe deformation typified by loss of metallographic features. Concomitantly, slip is expected to be the dominant deformation mode in this zone. Likewise, closest to the machined surface, large values of strain and a metallographically featureless microstructure are observed. This microstructure is likely to mirror that of the nanostructured chip. Further into the subsurface, the values of strain are smaller and significant twinning is observed. Acknowledgements We thank the US Department of Energy (Grant 4000031768 via UT-Batelle), Ford Motor Company, Oak Ridge National Laboratory (ORNL), the NSF (Grant DMI 0500216) and the State of Indiana 21st Century Research and Technology Fund for support of this work. Additional thanks are also due to Drs. Ray Johnson (ORNL) and Andrew Sherman (Ford) for their encouragement of the studies. References [1] Ferrasse S, Segal VM, Hartwig KT, Goforth RE. J Mater Res 1997;12:1253–61. [2] Stolyarov VV, Zhu YT, Lowe TC, Valiev RZ. J Nanosci Nanotechnol 2001;1:237–42. [3] Stolyarov VV, Zhu YT, Alexandarov IV, Lowe TC, Valiev RZ. Mater Sci Eng A 2003;343:43–50. [4] Stolyarov VV, Zhu YT, Alexandarov IV, Lowe TC, Valiev RZ. Mater Sci Eng A 2001;299:59–67. [5] Zhu YT, Huang JY, Gubicza J, Ungar T, Wang YM, Ma E, et al. J Mater Res 2003;18:1908–17. [6] Kim I, Kim J, Shin DH, Liao XZ, Zhu YT. Scripta Mater 2003;48: 813–7.
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