MICROSCOPY RESEARCH AND TECHNIQUE 69:374–381 (2006)
Scanning Electron Microscopy Imaging of Dislocations in Bulk Materials, Using Electron Channeling Contrast MARTIN A. CRIMP* Department of Chemical Engineering and Materials Science, Michigan State University, East Lansing, Michigan 48823
KEY WORDS
SEM; ECCI; twins; dislocation density; contrast analysis
ABSTRACT The imaging and characterization of dislocations is commonly carried out by thin foil transmission electron microscopy (TEM) using diffraction contrast imaging. However, the thin foil approach is limited by difficult sample preparation, thin foil artifacts, relatively small viewable areas, and constraints on carrying out in situ studies. Electron channeling imaging of electron channeling contrast imaging (ECCI) offers an alternative approach for imaging crystalline defects, including dislocations. Because ECCI is carried out with field emission gun scanning electron microscope (FEG-SEM) using bulk specimens, many of the limitations of TEM thin foil analysis are overcome. This paper outlines the development of electron channeling patterns and channeling imaging to the current state of the art. The experimental parameters and set up necessary to carry out routine channeling imaging are reviewed. A number of examples that illustrate some of the advantages of ECCI over thin foil TEM are presented along with a discussion of some of the limitations on carrying out channeling contrast analysis of defect structures. Microsc. Res. Tech. 69:374– 381, 2006. V 2006 Wiley-Liss, Inc. C
INTRODUCTION Since its first application to the study of crystalline materials nearly 50 years ago, amplitude contrast based TEM has been the most widespread technique for the study and analysis of dislocations (for a review see Hirsch et al., 1965). TEM has the advantage over many techniques in being able to assess the details of dislocation structures at a local level in relation to surrounding microstructural features. By using it in conjunction with selected area diffraction, details such as dislocation line directions and dissociation planes can be readily assessed. Diffraction contrast analysis, using different diffracting conditions to obtain different image characteristics, allows even more detailed analysis, such as dislocation Burgers vectors determination using g b ¼ 0 invisibility criteria. The primary drawback in carrying out the analysis of dislocations using TEM methods continues to be the need for extremely thin samples. Historically, preparation of TEM thin foils has been carried out using electropolishing or ion milling, and more recently, tripod polishing. These approaches have been very tedious, and for some materials, nearly impossible. With the advent of focused ion milling (FIB), TEM sample preparation has become more reliable, but with a significant capital outlay. Regardless, TEM thin foil analysis of dislocations still has a number of significant limitations that are inherent to the thin foil approach. Utmost is the question as to whether the sample preparation has influenced the number and morphology of dislocations. Ion milling (and FIB) can introduce point defects that in turn may alter the state of the dislocations. Bulk thinning prior to electropolishing, or the very nature of tripod polishing, may introduce dislocations through deformation. Furthermore, image forces, C V
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which may alter the dislocation distribution, may become severe in thin foils. Thin foil preparation also tends to result in limited observation area, leading to the question of whether the observations are isolated anomalies or representative of the broader material. Because it is often difficult to select the specific area where thinning will occur (with the exception of FIB), it is difficult to prepare thin foils that allow examination of critical volumes of materials, such as specific grain boundaries or crack tip regions. Finally, because of the very nature of TEM thin foils, carrying out in situ experiments can be difficult. For example, in situ TEM deformation experiments can be difficult to interpret because of the complex state of stresses associated with randomly shaped thinning holes and tapered foil edges. Consequently, most TEM analysis of dislocation structures in deformed materials is carried out postdeformation, making it impossible to accurately relate pre- and postdeformation morphologies in context with specific microstructural features. While diffraction contrast TEM has proven to be a very versatile tool for assessing dislocation structures and behavior, the limitations outlined above have led to a long series of attempts to develop a scanning elec*Correspondence to: Martin A. Crimp, Department of Chemical Engineering and Materials Science, Michigan State University, East Lansing, Michigan 48823, USA. E-mail:
[email protected] Received 2 September 2005; accepted in revised form16 September 2005 Contract grant sponsor: National Science Foundation; Contract grant number: DMR 9257826; Contract grant sponsor: The Office of Naval Research; Contract grant number: N00014-94-1-0204; Contract grant sponsor: The Air Force Office of Scientific Research; Contract grant number: F49620-01-1-0116; Contract grant sponsor: The MSU Composite Materials and Structures Center; Contract grant sponsor: The National Science Foundation; Contract grant number: DMR 9302040; Contract grant sponsor: Michigan State University. DOI 10.1002/jemt.20293 Published online in Wiley InterScience (www.interscience.wiley.com).
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Fig. 1. Schematic representation of the variation in back scattered electron emission (h) as a function of the relative orientation between the incident electron beam and the crystal lattice (after Joy et al. 1982).
tron microscopy based approach to carry out the type of analysis available in TEM. These approaches have been based on the phenomena of electron channeling discovered by Coates in 1967, who found a strong variation in the backscattered electron yield as function of the relative orientation between the incident electron beam and the crystal (Fig. 1) that resulted in bands of contrast that were proportional in width to the accelerating voltage (Joy et al., 1982). In comments on Coates’ work, Booker et al. (1967) outlined that the formation of channeling bands could be interpreted by the superposition of two Block wave functions. Consequently, Booker et al. theorized, ‘‘It should in principle be possible to use the scanning electron microscope to detect dislocations by the direct examination of unetched crystal surfaces. It is only necessary to orientate the crystal at the Bragg position and the local bending of the crystallographic planes should produce the necessary contrast. Such contrast should be directional and could lead to Burgers vector determination.’’ While Booker et al. outlined the general framework to carry out channeling contrast imaging of dislocations, they also outlined a number of potential difficulties in carrying out such analysis. Both Clarke and Howie (1971) and Spencer et al. (1972) carried out dynamical theory calculations of dislocations and stacking faults in thin and bulk specimens. It was found that because the backscattered electron (BSE) signal in bulk samples is dominated by the BSEs in the interaction volume rather than the electrons scattered from the primary beam, the contrast from dislocations in bulk samples is much lower than in thin foils. Taking in to account current density, beam divergence, and spot size, Spencer et al. concluded that under optimum conditions, dislocations might be observed in thin samples, using a conventional tungsten filament in an SEM. However, to observe dislocations in bulk samples, it would be necessary to use a field emission gun (FEG) electron source. Consequently, the first experimental images of disloca-
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tions using backscattered electrons were captured from TEM thin foils, using scanning TEM (STEM), by Clarke (1971) and Stern et al. (1972). The first successful attempt to image dislocations in bulk samples by using channeling contrast were carried out by Pitaval et al. (1977) and Morin et al. (1979). This group used a cold FEG SEM operated at 40–50 kV and tilted the sample 508–708 to increase the total number of BSEs and to optimize the number of high energy BSEs. In addition, the group used a high-energy filter located in front of a scintillator detector, to remove lower energy BSEs in order to increase the signal to noise ratio (SNR), as predicted by Sandsto¨m et al. Removal of these lower energy electrons effectively limited the collected electrons to a smaller interaction volume. With this experimental configuration, this group was able to demonstrate dislocation line resolutions of 10 nm with observation depths of 100 nm in Si. The group also showed that the g b ¼ 0 invisibility criteria could be applied to dislocations observed using channeling contrast. To overcome the low SNR while avoiding the difficulties of applying an energy filter, Czernuszka et al. used digital signal processing to collect channeling contrast images of dislocations in Si and Ni3Ga (Czernuszka et al., 1990a,b; 1991). Using a high tilt configuration in a cold FEG to enhance the BSE signal, this group demonstrated that dislocation images could be drawn out of the weak signal by carrying out a background subtraction and averaging over multiple image frames (10 frames with 10 s frame time). This approach greatly simplified carrying out what they termed as electron channeling contrast imaging (ECCI). Because no energy filter was used, this group was able to carry out experiments over a wide range of accelerating voltages (Czernuszka et al., 1991), finding that image contrast was enhanced at lower voltages (30 kV), while imaging depths were greater at higher accelerating voltages (100 kV), in agreement with the earlier theoretical work of Spencer et al. (1972). Using this approach, a number of studies have been carried out that used ECCI to examine dislocations. Wilkinson et al. (1993; 1994) used ECCI to show that accumulations of interfacial dislocations could be imaged to depths greater than 1 lm in Si Ge epilayers. However, at such depths the spatial resolution was found to be too low to image individual dislocations in the epilayers. Rather, groups of clustered dislocations were imaged. A number of studies have taken advantage of the relative ease of imaging accumulations of dislocations, using channeling contrast to study dislocation cell morphology in fatigued metals (Ahmed et al., 1997; 2001; Bretschneider et al., 1997; Chen et al., 1997; Melisova et al., 1997; Zauter et al., 1992; Zhai et al., 1996). Recently, Dudarev et al. (1999) developed a model for the channeling contrast from dislocations in cell walls. While these previous studies have demonstrated the ability to carry out imaging of dislocations in bulk studies, they have not achieved the goal of having a robust alternative to TEM for the routine imaging and analysis of individual dislocations. Over the past decade, significant progress towards this goal has been achieved (Ng et al., 1998; Simkin and Crimp, 1999). These studies have demonstrated that ECCI of individual disloca-
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tions can be carried out over a wide range of conditions, using a straightforward experimental configuration, with a standard commercially available scanning electron microscope. The present paper summarizes this experimental configuration and illustrates how this set up facilitates a wide range of experimental studies. A number of case studies that illustrate the unique capabilities of using ECCI for material analysis are presented. EXPERIMENTAL CONFIGURATION Studies have been carried out using a CamScan 44 FE SEM equipped with a Schottky thermal field emission gun operated at 25 kV. An advantage of this system over the cold field emitters typically used in the past is that the Schottky emitter has significantly higher emission currents, leading to greatly enhanced signal strengths, with only a small increase in probe size at the accelerating voltage used. The experimental electron beam parameters were as follows: probe current, 1.8–2.9 nA; beam convergence angle (2a), 5–7 mrad; and spot size, 20–30 nm. Working distances of 10–13 mm were used with sample tilts ranging from 08–308. The CamScan 44 FE was equipped with a manufacturer-supplied selected area electron channeling (SACP) module. SACPs allow crystallographic orientations to be determined, and in particular, specific channeling (Bragg) conditions to be established by setting up ‘‘two-beam’’ imaging conditions by aligning the untilted electron beam (optic axis of the microscope) to the edge of a channeling band. While some ECCI can be carried out without SACP capabilities (particularly for single crystals where channeling patterns can be observed at low magnification), the ability to collect orientation information from small areas (down to 10 lm) greatly increases the flexibility of the system. It should be noted that electron backscattered diffraction patterns (EBSD patterns) are not useful for establishing the channeling conditions for ECCI analysis, as EBSD determines the crystal orientation relative to a calibrated highly tilted specimen normal (with accuracy in the range of 18), while SCAP establishes the crystal orientation relative to the incident electron be trajectory (with an accuracy of a fraction of a degree). Although the setting up of two-beam channeling conditions is in many ways analogous to the setting up of two-beam diffracting conditions in TEM, there is one important difference. In TEM diffraction contrast imaging, it is customary to use a positive deviation from the Bragg condition (s > 0) to optimize image contrast (Hirsch et al., 1965). This condition is achieved by moving the bright Kikuchi line outside of the corresponding diffraction spot. However, in ECCI, both theoretical (Spencer et al., 1972) and experimental (Simkin and Crimp, 1999) studies have shown that maximum dislocation contrast occurs when an s ¼ 0 imaging condition is used. Such a condition is achieved by tilting the specimen so that a strong channeling band passes over the center point of the channeling pattern (as shown in Fig. 2), which results in the un-tilted beam (center of the pattern, which will be the beam trajectory during image rastering) being at the exact Bragg channeling condition. Past studies of dislocations using channeling contrast have taken advantage of the enhanced number of
Fig. 2. Set up of channeling conditions for ECCI. (a) Obtain a low camera length SACP. (b) Tilt sample to a prominent channeling orientation. (c) Using a high camera length, adjust sample tilt so that the edge of the channeling band lies at the optic axis of the microscope. This results in a s ¼ 0 condition during imaging.
BSEs at high tilt angles (Golstein et al., 1992) by using side mounted BSE detectors to collect the channeling signal from highly tilted samples. In contrast, our recent studies have used a standard 4-quadrant Si diode detector mounted on the pole piece of the objective lens to image samples at relatively small tilts. Figure 3 contrasts the two imaging configurations and shows dislocation images and SACPs collected under these two experimental arrangements. While the overall BSE signal must be lower in the low tilt configuration, it has been found that the large collection angle (0.6p str) offered by the pole piece-mounted Si detector coupled with the high brightness Schottky source
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Fig. 3. Two different imaging configurations for carrying out ECCI shown with images of dislocations in Si and associated SACPs. (a) Low tilt configuration with pole piece-mounted Si detector results in minimal foreshortening (c) and even SACP intensity (e). (b) High tilt configuration results in significant image foreshortening (d) and large intensity variations across SACP (f).
offer adequate signal to collect images using standard recursive filtering of 32 frames with a 0.85 s frame time. The image foreshortening that is a consequence of using a highly tilted sample is avoided (with the sample tiled 658, the image compression is 2.43 in the vertical direction in Fig. 3d). SACP interpretation is also greatly enhanced using a low tilt configuration, which results in consistent contrast across the pattern. In the high tilt configuration, the large changes in BSE signal with small variations in beam tilt (Fig. 3f) results in pattern saturation on one side of the pattern and signals below the detection limit on the other side of the pattern, making orientation determination and channeling band identification difficult (Simkin and Crimp, 1999). In addition to making the image and SACP interpretation more straight forward, the low tilt configuration also has the advantage of allowing a wider range of specimen tilts to be imaged, and allows greater flexibility in the microscope chamber. In our studies, this has allowed the use of a number of in situ deformation stages in conjunction with ECCI studies of dislocations. ECCI STUDIES OF DISLOCATION STRUCTURES Using the experimental configuration described earlier, near surface dislocations can be studied in bulk
samples, with a number of advantages over conventional TEM studies: 1. Ease of sample preparation. Only one surface of the sample needs to be prepared, rather than a thin TEM foil. Typically, this surface can be prepared and evaluated for sample preparation artifacts prior to further study. For example, the pre-existing dislocation structure can be evaluated prior to bulk specimen treatment such as deformation. Most TEM sample preparation is carried out after treatment, leading to uncertainty as to whether sample preparation has altered the dislocation structure. 2. Large examination area. Because bulk samples are used, the area available for examination is essentially limited to the entire specimen surface, rather than a limited thin area. This allows either specific areas of interest to be studied (for example specific interfaces or grain boundaries) or surveys across a sample that allows true representative distributions to be determined. 3. In situ studies. Because bulk samples are examined in ECCI, it is particularly well-suited for carrying out in situ studies, allowing the development of defect structures to be characterized in the same areas over time. Bulk samples allow for a well-characterized bulk state of stress to be correlated with the details of dislocation behavior at the crystal
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Fig. 4. (a) Secondary electron image of a crack tip in an NiAl single crystal deformed in 4-point bending. (b) Electron channeling contrast image of the same region showing dislocations emitted from the crack tip. Dislocations to the upper right of the crack tip appear as dots, because of their line directions being nearly normal to the free surface. Dislocations to the upper left of the crack tip appear more as lines, as they are parallel to the surface.
level. In contrast, in situ deformation TEM studies are plagued by the uncertainty of the complex state of stress that develops as a result of a TEM thinning hole and tapered thin sections. In the following sections, a number of example studies will be presented that illustrate some of the advantages of using ECCI for evaluating crystallographic defect, in particular, dislocations. Examination of Dislocations at Crack Tips and Edges The nature of dislocation generation in relation to crack propagation has been investigated in NiAl single crystals, using in situ deformed 4-point bend specimens (Ng et al., 1997, 1998). Dislocations have been examined in the near surface regions of the specimen face normal to the notch root. The specimen was electropolished prior to in situ deformation. Figure 4a shows a secondary electron image of a crack tip, with the bulk tensile stress direction on the surface (plane stress) noted. The image shows low contrast due to the flat-
Fig. 5. (a) Secondary and (b) channeling contrast image of the crack path 1 mm from the location of Figure 4. The crack has propagated from the lower left towards the upper right. Dislocations have been asymmetrically generated as the crack passed, with dislocations arranged in dense slip bands on the lower side of the crack and more randomly distributed dislocations on the upper side of the crack.
ness of the electropolished surface, with only a small amount of topographical contrast associated with displacements along the crack. Figure 4b shows the same area imaged under optimum channeling conditions. The region above the crack tip shows small black dots, which are dislocations with their line directions near perpendicular to the surface (close to parallel to the crack front). White dots below the crack tip are also dislocations imaged under slightly different channeling conditions (some lattice rotation occurred around the crack). The region immediately along the crack appears bright due to backscattered electrons escaping from inside the crack as the electron beam was rastered near the crack edge. The size of the plastic zone in which the crack arrested can be readily measured, in this case, is in the range of 10 lm. Because imaging is carried out on a bulk sample, it is also possible to examine the dislocation structure at any point along the crack path, allowing the nature of plastic zone as the crack propagated to be investigated. Figure 5 shows a secondary electron/channeling image pair of crack path taken 1 mm from the tip shown in Figure 4. Along the crack path, well-defined slip bands are evident on one side of the crack, while on the other
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Fig. 6. Electron channeling contrast image of deformation twins and dislocations in g-TiAl.
side of the crack dislocations are in more random orientations, some with line directions near parallel with the surface. Because these studies are carried out in situ on bulk samples, it is possible to continue the loading to determine how the dislocation structure develops and the crack propagates after the initial crack arrest. In all of the cases observed, it was found that dislocations that developed during the initial loading did not move during reloading, and additional crack propagation did not occur from the arrested crack tip, but instead a new crack trip developed at some point along the existing crack, behind the arrested crack tip (Ng et al., 1997). Deformation Strain Transfer at Grain Boundaries Because ECCI is carried out on bulk samples, it is ideally suited for examining the nature of deformation mechanism interaction at grain boundaries. This has been studied in g-TiAl by examining the tensile surface of ex-situ deformed smooth 4-point bend samples that had been deformed to 1% plastic strain (Simkin et al., 2003a,b; 2002). Both dislocations and deformation twins are readily imaged in TiAl using ECCI (Fig. 6). In this material, grain boundary microcracks have been observed to develop at the intersection of deformation twins with g-g grain boundaries, an example of which is shown in Figure 7. The orientations of individual grains forming these microcracked boundaries have been determined using a combination of EBSD and SACP (Simkin et al., 2003a) (because the tetragonality of the L10 TiAl is small, EBSD cannot alone delineate between the c and a axes of the unit cell). Based on this orientation information, trace analysis has allowed unambiguous identification of the 4 {111} twining planes. As there is only one possible a/6 twinning direction on each {111} plane in TiAl, this orientation/trace analysis has allowed complete identification of the twinning systems involved in microcrack nucleation. Furthermore, because bulk samples have been used, the direction of the global tensile stress is known, allowing calculation of the Schmid factors for all of the possible twinning and dislocation slip sys-
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Fig. 7. Electron channeling contrast image of deformation twins interacting with a g-g grain boundary in TiAl. SACP based trace analysis allows identification of the active twinning planes. Grain boundary microcracks (arrows) have developed in association with the interaction of {111} deformation twins in the upper grain with the grain boundary.
tems, and the absolute directions of the twinning shears (unlike dislocation slip, twinning in unidirectional). Examination of both microcracked and uncracked boundaries has allowed development of a fracture initiation parameter that correlates with the propensity for a given boundary to develop microcracks, based upon the activity of twinning and dislocation slip systems and the misorientation of the deformation systems across the grain boundary (Simkin et al., 2003). Dislocation Density Measurements in Commercial Purity Ti Unlike TEM, where dislocation density measurements are limited to the thin section of a foil sample, ECCI readily allows dislocation densities to be measured over a very large area of a prepared surface. Consequently, ECCI is ideally suited for carrying out rapid dislocation density measurements from widely spaced representative areas. Dislocations have been imaged in commercial purity Ti, cold rolled to thickness reductions of 1.08, 2.16, 5.17, and 15.05% (Crimp et al., 2004). In the 1.08 and 2.16% deformed materials, individual dislocations were readily observed, with distinct slip bands evident in the 2.16% deformed Ti (Fig. 8a). At greater deformations, it was more difficult to resolve individual dislocations, with higher density volumes displaying mottled contrast, as shown in Figure 8b. Such contrast is not unlike TEM images of volumes with large dislocation densities. Dislocation densities were determined by simply measuring the number of dislocation/surface intersections within a defined area. A correction factor of 2 was applied to account for random dislocation orientations (Crimp et al., 2004). Densities of 7.8 3 1012, 6.8 3 1013, and 1.2 3 1014 m2 were measured for the 1.08, 2.16, and 5.17% deformed materials, respectively. Because of inability to consistently resolve individual dislocations in the 5.17% deformed Ti, 1.2 3 1014 m2 should be considered a lower bound for the dislocation density at
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While both the g b ¼ 0 (Morin et al., 1979; Czernuszka et al., 1990b; Simkin and Crimp, 1999) and the g b 3 u ¼ 0 (Crimp et al., 2001) invisibility criteria have been demonstrated for ECCI, in practice, significant care must be taken when interpreting such contract analysis. TEM allows reasonable images of dislocations to be obtained for a significant range of positive deviations from the Bragg condition (Hirsch et al., 1965). In contrast to TEM, electron channeling imaging of dislocations is optimized with s ¼ 0, and the image contrast falls off rapidly with both positive and negative deviations from the perfect Bragg condition (Spencer et al., 1972; Simkin and Crimp, 1999). Furthermore, collection of channeling images of dislocations requires careful attention to focus (as opposed to TEM where the depth of focus for objects like dislocations is very large compared to the sample thickness) and proper background subtraction and frame averaging. Because of these reasons, it is quite possible to obtain images with little or no dislocation contrast under channeling conditions that should produce dislocation contrast. Consequently, care must be taken not to wrongly interpret a lack of dislocation visibility as an invisibility condition. To properly interpret channeling contrast invisibility, one should ensure that other features, such as dislocations with different Burger’s vectors, are in strong contrast. Furthermore, interpreting channeling contrast images with additional information such as plane trace analysis and even complementary TEM analysis can assist in proper ECCI analysis. ACKNOWLEDGMENTS The author acknowledges the assistance of B.A. Simkin, B.C. Ng, and J. Hile in carrying out the ECCI studies summarized in this article, and also the helpful discussions with T.R. Bieler. Fig. 8. Electron channeling contrast imaging of dislocations in commercial purity Ti. (a) Distinct slip bands running from the lower left towards the upper right are apparent in the 2.16% cold rolled material. (b) Higher dislocation densities in the 5.17% cold rolled material make it more difficult to resolve individual dislocations, resulting in mottled contrast.
the deformation level. More importantly, however, this sets an approximate dislocation density of 1014 m2 as an upper limit for using ECCI for imaging dislocations. This limit is very similar to the limit of 3 3 1013 m2 given by Hirsch et al. (1965) for measuring dislocation densities in TEM by counting dislocation interactions with the two free surfaces in a TEM thin foil (note that the approach outlined by Hirsch et al. does not multiply the measured dislocation density by 2 to take in to account the random orientations of the dislocations). ECCI CONTRAST ANALYSIS AND LIMITATIONS As demonstrated above, because ECCI is not limited to thin foil samples, it has the potential to allow characterization of dislocation structure in relation to specific microstructures and samples geometries that are not possible with TEM. However, ECCI is not without limitations.
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